Chemical Solution Deposition of Semiconductor Films

Marcel Dekker, Inc. New York • BaselTM

CHEMICAL SOLUTIONDEPOSITION OF

SEMICONDUCTOR FILMS

Gary HodesWeizmann Institute of Science

Rehovot, Israel

Copyright © 2002 by Marcel Dekker, Inc. All Rights Reserved.

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To my parents, for their dedication to my educationand their faith in my abilities


Preface

Chemical solution deposition (CSD); known also as chemical bath deposition(CBD) and simply chemical deposition (CD, the form we will use in this book)was first described in 1869, and it has been used since to deposit films of manydifferent semiconductors. It is probably the simplest method available for this pur-pose—all that is needed is a vessel to contain the solution (an aqueous solutionmade up of a few, usually common, chemicals) and the substrate on which depo-sition is required. Various “complications,” such as some mechanism for stirringand a thermostated bath to maintain a specific and constant temperature, are op-tions that may be useful.

In spite of this extreme experimental simplicity, understanding the mecha-nisms involved in the deposition and the ability to widen the range of deposits ob-tained—both in composition and the control of numerous other properties—isusually not so simple. Also in spite of its simplicity, it has not been exploited as atechnique as much as might be expected. However, CD has experienced some-what of a renaissance recently, due largely to its overwhelmingly successful usein depositing buffer layers of CdS (and similar materials) in thin-film photovoltaiccells. The deposition of the CdS, as with many other semiconductors that havebeen deposited by CD, is often recipe oriented; there seem to be almost as manydifferent “recipes” as there are groups.

Notwithstanding the wide interest and use of this technique, at the timewhen the idea to write this book was conceived, there was no recent comprehen-


sive review or any general summing up of the field of CD. The first general re-view on the subject is that by Chopra et al. in 1982 [1]. Nine years later, a reviewby Lokhande [2] was published with an emphasis on describing the deposition ofthe various semiconductors that had been deposited up to then. A comprehensiveand general review was published by Lincot et al. in 1998 [3] (about a year afterthe writing of this book had commenced) and, in the same year, two more reviews,one a more specific review by Nair et al. [4] on their extensive work in the fieldconnected with solar energy–related issues and the other, by Savadogo [5], de-scribing CD (and electrodeposited) semiconductors used as solar energy materi-als. In the last two years, two new reviews have appeared, by Mane and Lokhande[6] and most recently by Niesen and DeGuire [7], the latter covering also other so-lution deposition methods, such as SILAR (Successive Ion Layer Adsorption andReaction), electroless deposition and liquid phase deposition (which can be con-sidered to be a subset of CD) and emphasizes oxides, although sulphides and se-lenides are also covered.

The driving force for this book is the perceived need for a detailed coverageof a field that has expanded enormously in recent years. While the title of the tech-nique suggests that the book is aimed mainly at chemists, this would be an incor-rect impression. Many of those who should find this book useful will be “physi-cists” or “engineers” who are dealing with thin-film photovoltaic cells. Some ofthese readers may have only a superficial background in chemistry, and for these,Chapter 1, Fundamentals, which deals with the science (largely chemistry) behindthe technique, will be very important background reading. In this chapter, mate-rial will be found on such topics relevant to CD as principles of precipitation andsolubility product; nucleation; growth; colloids; aggregation and sticking. Eventhose with a good chemistry background are advised to read this chapter, if onlyto refresh their knowledge.

In one respect, this book is organized somewhat differently than usual. Itcontains a fairly comprehensive review of CD in the form of Chapter 2, GeneralReview. Most of the material in this review will be expanded on in the relevantchapters, and one might ask why it is included at all. The reason is that most peo-ple do not read a book of this type from cover to cover; they read those chaptersor parts of chapters they consider relevant to their purposes. In doing this, they arelikely to miss matter from other chapters that is also relevant. For this reason, it isstrongly recommended that all readers read through the first three chapters. Chap-ter 1, as already noted, is to acquire or refresh the relevant scientific background.Chapter 2 should give a good overview of what has been done without having togo into too much detail. Apart from this, however, Chapter 2 contains detail notfound elsewhere; a short history of CD, details on the effects of substrate natureand variation (if I were to rewrite this book, an additional chapter would be de-voted solely to the substrate), some “recipes” for depositing certain films and,very important, descriptions, and, where possible, explanations of the reasons for


a particular recipe and the expected effects of various changes to those recipes.Also, other methods that are related to CD are treated briefly in this chapter. In or-der to promote the ease of reading and “flow” in this chapter, references are, forthe most part, not given unless they are not clearly provided in the chapter and sec-tion relevant to that particular topic elsewhere in the book. The importance of thethird “required” chapter, Chapter 3, Mechanisms of Chemical Deposition, shouldbe self-evident. Understanding the possible mechanisms of the deposition is thebest defense against what is commonly (and often with considerable justification)thought of as a “recipe-oriented” field. Not that that a reading of this chapter willautomatically endow the reader with the ability to know what mechanism is actu-ally operating in every case—unfortunately, often far from it. However, it is hopedthat it will help in the movement toward that goal.

Having worked (it is hoped) through these “compulsory” chapters, mostreaders will want more detail on specific aspects that are important to them. Howcan films of X be made? How are the various experimental parameters expectedto affect the properties of this film? Why is CD so useful for photovoltaic cells?How can nanocrystalline films of Y be made with a specific crystal size? Suchquestions will be answered in subsequent chapters (or, if not answered, at least in-formation will be given to allow the reader to plan experiments in order to find theanswers).

The next five chapters deal with deposition of specific groups of semicon-ductors. In Chapter 4, II–VI Semiconductors, all the sulphides, selenides, and(what little there is on) tellurides of cadmium (most of the chapter), zinc (a sub-stantial part), and mercury (a small part). (Oxides are left to a later chapter.) Thischapter is, understandably, a large one, due mainly to the large amount of workcarried out on CdS and to a lesser extent on CdSe. Chapter 5, PbS and PbSe, pro-vides a separate forum for PbS and PbSe, which provided much of the focus forCD in earlier years. The remaining sulphides and selenides are covered in Chap-ter 6, Other Sulphides and Selenides. There are many of these compounds, thus,this is a correspondingly large chapter. Chapter 7, Oxides and Other Semicon-ductors, is devoted mainly to oxides and some hydroxides, as well as to miscella-neous semiconductors that have only been scantily studied (elemental seleniumand silver halides). These previous chapters have been limited to binary semicon-ductors, made up of two elements (with the exception of elemental Se). Chapter8, Ternary Semiconductors, extends this list to semiconductors composed of threeelements, whether two different metals (most of the studies) or two differentchalcogens.

The final two chapters deal with “applications” (in the scientific as well ascommercial sense) of CD films. As already mentioned, photovoltaic cells is theone subject that has given CD a push in the last decade, while photoelectrochem-ical cells was probably the main driving force for such studies in the decade be-fore that. Chapter 9 deals with Photovoltaic and Photoelectrochemical Properties.


Finally, the tendency for CD films to be nanocrystalline and often to exhibit quan-tum-size effects is treated in the final chapter, Chapter 10, Nanocrystallinity andSize Quantization in CD Semiconductor Films.

The layout of this book means that there will be some overlap between sec-tions. However, this system should allow those readers interested in one or morespecific sections to skip the others, thereby making the book more efficient for theindividual reader. An example of this is the use of quantum-size effects to eluci-date CD mechanisms. This is treated, with different emphasis, both in Chapter 3(Mechanisms of Chemical Deposition) and in Chapter 10 (Nanocrystallinity andSize Quantization in CD Semiconductor Films).

This book covers a field that, from its title, may appear to be limited. In re-ality, it is surprisingly multidisciplinary. Inorganic chemistry and film formationare, of course, fields that are evident from the title. Certainly, those concernedwith the deposition of semiconductor films for any purpose should find this bookuseful and informative. However, it will also be valuable to those working in otherfields. The considerable section on semiconductor quantum dots, for example,will be of interest to those working with low-dimensional semiconductors. Scien-tists and engineers working in thin-film solar cells will find a compendium of re-search on CD buffer layers in these cells. There is much in the book relevant tocolloid scientists. Even biologists are not forgotten: the slow formation and depo-sition of inorganic compounds characteristic of CD has a lot in common withbiomineralization, and hopefully this book will be useful to those working in thefield of biomimetics. And, of course, the original application of CD, for near- tomid-infrared detectors, will attract those designing or using optoelectronic equip-ment in this wavelength range.

Finally, a word concerning the coverage of the literature. When starting outon this enterprise, the intention was to try and cover the field more or less com-pletely, with the exception of some of the earlier work, mostly on PbS. Of course,some papers might occupy a considerable amount of book space, while, at theother extreme, others might just be mentioned in a table. During the long courseof putting this book together, it became increasingly clear, from the “new” litera-ture that appeared (not necessarily chronologically new but just new to the au-thor), even toward the final stages of writing, that an appreciable number of rele-vant papers would remain unknown to the author. For these, the author expressesregret, not only to the authors of such works, but also to the readers of this book.It is hoped that the expanding literature that is continually appearing on the sub-ject does not render this book out of date too rapidly. From the objective scientificviewpoint, it can be hoped that this process will not be too slow either.


REFERENCES1. KL Chopra, RC Kainthla, DK Pandya, AP Thakoor. In: Physics of Thin Films, Vol. 12.

Academic Press, New York and London, 1982, pp 167.2. CD Lokhande. Mater. Chem. Phys. 28:1, 1991.3. D Lincot, M Froment, H Cachet. In: RC Alkire, DM Kolb, Eds. Adv. Electrochem. Sci.

Eng., New York: Wiley-VCH, 1998, Vol. 6, p 165.4. PK Nair, MTS Nair, VM Garcia, OL Arenas, Y Pena, A Castillo, IT Ayala, O Gomez-

daza, A Sanchez, J Campos, H Hu, R Suarez, ME Rincon. Sol. Energy Mater. Sol. Cells52:313, 1998.

5. O Savadogo. Sol. Energy Mater. Sol. Cells 52:361, 1998.6. RS Mane, CD Lokhande Mater. Chem. Phys. 65:1, 2000.7. TP Niesen, MR De Guire J. Electroceram. 6:169, 2001.


Contents

DedicationPreface

1. Fundamentals

2. General Review

3. Mechanisms of Chemical Deposition

4. II-VI Semiconductors

5. PbS and PbSe

6. Other Sulphides and Selenides

7. Oxides and Other Semiconductors


8. Ternary Semiconductors

9. Photovoltaic and Photoelectrochemical Properties

10. Nanocrystallinity and Size Quantization inChemical Deposited Semiconductor Films


1Fundamentals

The purpose of this chapter is to give sufficient background in the chemical prin-ciples involved in CD. For those with a good background in chemistry, a quick readthrough this chapter as a refresher course will probably be sufficient. However, thechapter is written, to a large extent, keeping in mind that not all readers will havea good chemistry background. The emphasis is on a qualitative or semiquantitativeunderstanding of the principles involved, sufficient to understand the concepts thatarise throughout the book. A deliberate policy has been made not to go too deeplyinto the details of these fundamentals where it is considered unnecessary; refer-ences to further reading will be given for those who wish such additional detail.

1.1 SOLUTION CHEMISTRY

1.1.1 Basic Terminology

The pH of a solution is the negative logarithm of the hydrogen ion concentrationin the solution:

pH � �log [H�] (1.1)

(Note for thermodynamic purists: Here and for the rest of this book, concentra-tions are used in place of activities.) Concentrations are denoted by square brack-ets; thus [H�] means the concentration of hydrogen ions.


The pH of pure water at 25°C is 7. Most (but not all) CD reactions take placein basic solutions at typical pH values of 9–12. Since hydroxide intermediates areoften important in CD, it is worth noting that a pH of 10 is equivalent to a hy-droxide ion concentration of 10�4 M at 25°C (since the ion product of water,[H�][OH�], � 10�14 at this temperature). As will be discussed shortly, this ionproduct is very temperature dependent, and so the OH� concentration at any par-ticular pH varies considerably with temperature.

The pH of pure (and also not so pure) water is very sensitive to small con-centrations of acids and bases. One drop of concentrated sulphuric acid added toa liter of water will change the pH by 4 pH units (from 7 to ca. 3). Solution pH canbe stabilized by a buffer (although there may be cases where a stable pH is not de-sirable); addition of (not too large) quantities of acid or base to a buffered solutionwill not affect the pH much. Buffers are usually mixtures of weak acids or basesand their salts. A common example in CD is the use of an ammonium salt(NH4

�X�) to control the pH of an ammonia solution. The equilibrium of ammo-nia in water is given by

NH3 � H2OD NH4� � OH� (1.2)

Since hydroxide ions are formed when ammonia dissolves in water, the pHof an aqueous ammonia solution is alkaline. The value of pH can be calculatedfrom a knowledge of the equilibrium constant, K, of this equilibrium. The equi-librium constant for an equilibrium of general type

aA � bB � �� dDD eE � f F � �� hH (1.3)

is given by

K � (1.4)

K for the ammonia dissolution, Eq. (1.2), is given by

K � � 1.8 � 10�5 (at 25°C) (1.5)

For example, for a 1 M solution of ammonia ([NH3] � 1), since the NH4� and

OH� concentrations are equal [from Eq. (1.2)], [OH�] can be calculated to be4.2 � 10-3 M. Knowing that the ion product of water, [H�][OH�], � 10�14 at thistemperature allows us to convert [OH�] to [H�] and thus to find the pH (which is11.62).

If an ammonium salt is added to the ammonia solution, the NH4� concentra-

tion is now dictated by the concentration of ammonium salt added rather than bythat existing due to the weak dissociation of ammonia. Thus, assuming the same1 M ammonia as before, but adding 0.1 M NH4

� (say, as NH4Cl), then, ignoringthe few percent correction due to the extra NH4

� arising from the NH3 dissocia-

[NH4�][OH�]

��[NH3]

[E]e[F]ƒ . . . [H]h

��[A]a[B]b . . . [D]d


tion, [OH�] is given by

[OH�] � � 1.8 � 10�4 M (1.6)

and the pH of the solution becomes 10.25.The buffering action of this solution can be understood by considering equi-

librium Eq. (1.2). If extra OH� is added to the solution, the equilibrium is shiftedto the left; i.e., it tends to remove OH�. It also removes NH4

�; but if the concen-tration of this ion is high to begin with, then this change will not affect the pHgreatly. (This is the reason that the combination of ammonia and ammonium ionsis a better buffer than ammonia by itself.) A similar argument can be made for ad-dition of acid, through the equilibrium:

NH3 � H�D NH4� (1.7)

In this case, hydrogen ions are consumed in converting ammonia to NH4�.

This buffering action requires nonionized base (or acid) to operate, hencethe requirement of a weak base or acid together with its salt

As well as this buffering action, addition of ammonium ion also decreasesthe pH of an ammonia solution as shown above. This is an important effect—moreimportant than the buffering action in many CD processes.

1.1.2 Hydrolysis of Metal IonsMost cations are hydrated in aqueous solutions to a greater or lesser extent:

Mx� � nH2ODM(H2O)nx� (1.8)

The water is polarized and attracted by the positively charged cation. The greaterthe positive charge on the cation and the nearer the water can approach the cation,the greater will be this polarization and attraction. Thus small, highly charged(high-valence) cations will in general be more strongly solvated than large, mono-valent ones.

Continuing the argument, the positive charge on the cation attracts electronsfrom the oxygen of the water molecules. This, in turn, can result in the transfer ofelectron density from the OMH bonds to the (now electron-deficient) oxygen, asexemplified here:

(1.9)

This will weaken the OMH bond and may even break it, resulting in formation ofa metal hydroxide and a hydrogen ion, the latter which will be hydrated by a

1.8 � 10-5

��0.1


molecule of the surrounding water:

M(H2O)nx� � H2ODM(H2O)n�1OH(x�1)� � H3O� (1.10)

The H3O� (a hydrated hydrogen ion) is acidic; therefore this equilibrium gener-ates acidity in the solution. The more the cation attracts electron density from thewater, therefore, the more acidic is the cation. As with solvation, small, highlycharged cations should be more acidic than large cations with a small charge. Thecharge in particular is a very important factor in determining the degree of acidityof cations. Therefore monovalent cations are generally basic, while trivalent onesare acidic. Tetravalent cations, such as Sn4� and particularly Ti4� are so highlyacidic that their simple cations either do not exist in water or do so only under veryacidic conditions.

The hydrated metal hydroxy complex in Eq. (1.10) is a soluble species.However, if the pH is sufficiently high, the metal hydroxide, which is relativelyinsoluble for most metals (apart from the alkali group metals) will precipitate. ThepH value at which hydroxide precipitation occurs can be related to the acidity ofthe cation and is approximately equal to the pKa of the cation, where the pKa isminus the logarithm of the equilibrium constant of Eq. (1.10).

1.1.3 Solubility Product

A central concept necessary to understanding the mechanisms of CD is that of thesolubility product (Ksp). The solubility product gives the solubility of a sparinglysoluble ionic salt (this includes salts normally termed “insoluble”). Consider avery sparingly soluble salt (say, CdS) in equilibrium with its saturated aqueous so-lution:

CdS(s)D Cd2� � S2� (1.11)

(where subscript s represents the solid phase). The CdS dissolves in water to givea small concentration of Cd and S ions. This concentration is defined by the solu-bility product, Ksp, the product of the concentrations of the dissolved ions:

Ksp � [Cd2�][S2�] (1.12)

or more generally, for the dissolution:

MaXbD aMn� � bXm� (1.13)

Ksp � [Mn�]a[Xm�]b (1.14)

The more soluble is the salt, the greater the ion product and the greater is Ksp.However, Ksp also depends on the number of ions involved. Thus Bi2S3 has avalue of Ksp � [Bi3�]2[S2�]3 � 10�100. The very low value is due, in part, to therelatively large number of atoms in the Bi2S3 molecule and therefore of ions in-volved in the equilibrium. A list of approximate values of Ksp for some of the


semiconductors and related salts encountered in CD is given in Table 1.1. The val-ues for oxides are not so readily available as for sulphides and selenides. However,it must also be kept in mind that deposition of oxides often occurs via a hydrox-ide or hydrated oxide, and therefore the relevant value is that of this hydroxide orhydrated oxide.

Some explanation is required here concerning the S2� ion. In actual fact, inall but highly alkaline solutions (and the solutions used in CD, while mostly alka-line, are not that alkaline), most of the sulphur ion will be in the form of HS�

rather than S2�. This is due to the equilibrium between the two species:

HS�D S2� � H� Ka � 10�17.3 (1.15)

or alternatively, in terms of hydroxide concentration which is related to the hy-drogen ion concentration through the ion product of water:

HS� � OH�D S2� � H2O Ka � 10�3.3 (1.16)

(at room temperature).Thus at a pH of 11 (a common value in CD), which gives a value for [OH�]

at room temperature of 10�3 M, the S2� concentration will be

[S2�] � 10�3.3 [HS�][OH�] � 10�6.3 [HS�] (1.17)

Therefore the main sulphur ion in solution will be HS�.

TABLE 1.1 Values of the Solubility Product (at 25�C) for Compounds Relevant to CD

Solid Ksp Solid Ksp Solid Ksp Solid Ksp

Ag2S 3 � 10�50 CdSe 4 � 10�35 FeS 10�18 Pb(OH)2a 10�15–10�20

Ag2Se 10�54 CdTe 10�42 HgS 6 � 10�53 PbS 10�28

AgCl 2 � 10�10 Co(OH)2 5 � 10�15 HgSe 4 � 10�59 PbSe 10�37

AgBr 8 � 10�13 CoS 10�21 In(OH)3 6 � 10�34 Sn(OH)2 5 � 10�28

AgI 10�16 CuOH 1 � 10�14 In2S3 6 � 10�76 Sn(OH)4 1 � 10�56

As2S3 2 � 10�22 Cu(OH)2 2 � 10�20 Mn(OH)2 5 � 10�13 SnS 10�26

Bi(OH)3 6 � 10�31 Cu2S 10�48 MnS 10�13 SnS2 6 � 10�57

Bi2S3 10�100 CuS 5 � 10�36 Ni(OH)2 3 � 10�16 SnSe 5 � 10�34

Bi2Se3 10�130 CuSe 2 � 10�40 NiS 10�21 Zn(OH)2 10�16

Cd(OH)2 2 � 10�14 Fe(OH)2 5 � 10�17 NiSe 2 � 10�26 ZnS 3 � 10�25

CdS 10�28 Fe(OH)3 3 � 10�39 PbCO3 10�13 ZnSe 10�27

There is often a large variation in values from source to source—in some cases, some orders of magnitude. Forthis reason, only one significant figure (at most) is given before the exponent. A table of solubility products formany sulphides based on a reevaluated value for the second dissociation constant of H2S is given in Ref. 1. Thevalues in that study are typically some orders of magnitude lower than the ones shown here.a This is probably a hydrated lead oxide rather than a simple hydroxide.


Since Ksp is given in terms of [S2�], we can write [S2�] in terms of [H�]using Eq. (1.15):

[S2�] � 10�17.3[HS�]/[H�] (1.18)

and derive a solubility product of a sulphide, MxSy, in terms of the dominant HS�

concentration:

Ksp � [M2y/x�]x[S2�]y � [M2y/x�]x(10�17.3[HS�]/[H�])y (1.19)

A list of such solubility products for metal sulphides, as well as updated conven-tional ones, has been given by Licht [1]. In this book, we will continue to use themore conventional solubility products, partly because they are more common andpartly because the relevant equilibria are less unwieldly to describe.

In applying the solubility product concept to CD, it is often useful to con-sider it in terms of what concentration of ions is required in solution before pre-cipitation occurs. Thus, for CdS, with a Ksp value of ca. 10�28 (from Table 1.1),

[Cd2�][S2�] � 10�28 M. (1.20)

While the concentration of each ion in this example will be equal when dissolu-tion of the solid is considered, for formation of the solid from the ions, they maybe completely different; it is the product of the concentrations that is important.Thus a solution 0.2 M in sulphide ion and 10�27 M in Cd2� will (in principle) pre-cipitate CdS (the ion product will be greater than Ksp) while 0.1 M sulphide andthe same 10�27 M of Cd2� will (only just) not.

Ksp can be derived theoretically from the free energies of formation of thespecies involved in the dissolution equilibria. Thus, for the equilibrium

MaXb(s)D aMc�(aq) � bXd�(aq) (1.21)

the free energy of the dissolution is given by

�G0 � a�G0(aMc�(aq)) � b �G0(Xd�(aq)) � �G0(MaXb(s)) (1.22)

And since

�G0 � �RT ln K (1.23)

then

ln Ksp � ��RGT

0

� (1.24)

Since Ksp is a thermodynamic quantity, the ion product that should result in pre-cipitation may not necessarily do so for kinetic reasons (hence the term used ear-lier to qualify precipitation: “in principle”). This would be a case of supersatura-tion. In practice, however, the solubility product does give a fairly good idea ofwhen precipitation will occur in most cases.


CD reactions sometimes proceed via a metal hydroxide intermediate; theconcentration of OH� ions in the solution is particularly important in such cases.Since almost all CD reactions are carried out in aqueous solutions, the pH of thedeposition solution will give this concentration. In translating pH into OH� con-centration, the very temperature-dependent ionization constant of water should bekept in mind, as mentioned previously. The reason for this can be seen from Table1.2, which gives the OH� concentration in water at a pH of 10 (a typical pH valuefor many CD reactions), calculated from the ionization constant of water, Kw fromthe relation

�log Kw � �log [H�][OH�] � �log [H�] � log [OH�]

� pH � log [OH�](1.25)

The OH� concentration increases by nearly two orders of magnitude between 0and 60°C.

The OH� concentration increases (decreases) by one order of magnitude forevery unit increase (decrease) in pH. This means that the formation of a metal hydroxide (whether as a colloid or as a precipitate) in aqueous solution will bestrongly dependent on temperature when the product of the free metal ions andOH� ions is close to the hydroxide solubility product, although increase in Ksp

with temperature may partially offset this effect.

1.1.4 Complexation

Most CD reactions are carried out in alkaline solution. To prevent precipitation ofmetal hydroxides, a complexing agent (often called a ligand, since complexingagents to cations are electron donors) is added. The complexant also reduces theconcentration of free metal ions, which helps to prevent rapid bulk precipitationof the desired product. This section gives the basics of the theory of complexation.

TABLE 1.2 Effect of Temperature on OH� Concentrationin Water

Temp. OH� concentration(�C) -log10 Kw at pH � 10

0 14.944 1.138 � 10�5

10 14.535 2.917 � 10�5

20 14.167 6.808 � 10�5

30 13.833 1.469 � 10�4

40 13.535 2.917 � 10�4

50 13.262 5.470 � 10�4

60 13.017 9.616 � 10�4


If a KOH solution is added to a solution of a Cd salt, Cd(OH)2 will precipi-tate immediately. From the Ksp of Cd(OH)2 (2 � 10�14 at room temperature), andassuming a pH of, say, 11 ([OH]� � 10�3 at the same temperature), from Eq.(1.14), we can calculate that a Cd concentration above 2 � 10�8 M is enough toinitiate Cd(OH)2 formation.

If ammonium hydroxide (ammonia in water)—a common complexant forCd in CD—is added to a suspension of Cd(OH)2, the Cd(OH)2 will redissolve, as-suming enough ammonia has been added. How much is enough ammonia? Thiscan be calculated from the stability constant of the complex between ammonia andCd. The equilibrium of this reaction to form the cadmium tetraamine complex isgiven by

Cd2� � 4NH3D Cd(NH3)42� (1.26)

and the stability constant of this equilibrium, Ks, by

Ks � � 1.3 � 107 (1.27)

As calculated previously, Cd(OH)2 will precipitate when the free Cd2� concen-tration is larger than 2 � 10�8 M (at a pH of 11 and at room temperature). FromEq. (1.27), for a total Cd concentration of 0.1 M, we can calculate that a free NH3

concentration of 0.79 M will result in a free Cd2� concentration of 2 � 10�8 M.Add to this the ammonia tied up by complexation (0.1 � 4 M), the minimum NH3

concentration required to prevent precipitation of Cd(OH)2 is therefore 1.19 M. Ata Cd concentration of 0.01 M (more typical of many depositions), the corre-sponding concentration of ammonia is ca. 0.5 M. At a deposition temperature of60°C (CdS deposition is generally carried out at elevated temperatures, usually60°C), the ion product of water is 13, and therefore the OH� concentration at apH � 11 will be 10�2 M. Calculating the minimum concentration of ammonia re-quired to prevent precipitation of Cd(OH)2 at 60°C and 0.01 M total Cd gives avalue of 1.44 M. The value of pH chosen is typical of these solutions. For highervalues of pH, and at higher temperatures at the same pH (both of which mean anincreased [OH�]), more ammonia will be required to prevent precipitation ofCd(OH)2. This calculation ignored the decrease in the stability constant of thecomplex with increasing temperature (see later) as well as the increase in Ksp thatnormally occurs with an increase in temperature. These two effects act in oppositedirections; for most cases, their combined effect will be much smaller than that ofthe temperature-dependent ion product of water. Another simplification is the as-sumption of only one complex species; this simplification is reasonable for mostpurposes.

If a solution contains an excess of one of the ions of a sparingly soluble salt,this will modify the solubility of the sparingly soluble salt according to the com-mon ion effect. As an example of this effect, we might consider the precipitation

[Cd(NH3)42�]

��[Cd2�][NH3]4


of Zn(OH)2 by hydroxide according to the reaction

Zn2� � 2OH�D Zn(OH)2 (1.28)

An excess of OH� (the common ion) should shift the reaction to the right, i.e., tomore complete precipitation of the Zn(OH)2. This effect is a general one, but theconclusions are not always valid; the example (deliberately) given here is onewhere it is not valid. The reason is that OH� can form a complex with Zn2�

(Zn(OH)42�—the zincate ion), thus removing free Zn2� from solution and reduc-

ing the degree of precipitation. For a sufficiently high concentration of OH�,which can be calculated from the stability constant of the zinc–hydroxide (zincate)complex, the Zn(OH)2 will completely redissolve.

The stability constant of a complex is temperature dependent—increasedtemperature generally leads to increased dissociation of the complex. Qualita-tively, this can be explained by the Le Chatelier principle, which states that if thereis a change in a reaction parameter, the reaction will proceed in a direction that op-poses that change. Thus an increase in temperature will cause the reaction to go inthe direction in which heat is absorbed, which is dissociation of the complex.More quantitatively, the relation between equilibrium constant and temperature isgiven approximately by

ln K � ��

R�

TH0� (� a constant) (1.29)

where �H0 is the standard enthalpy change in the process and R is the gas con-stant. This is the integration of the van’t Hoff equation, hence the constant term.(The derivation of this equation can be found in any elementary physical chem-istry textbook and there is no need to repeat it here—the result is what is impor-tant for us.) This equation is approximate for a number of reasons. One is that itignores changes in entropy that often will act in the opposite direction for complexformation. However, the trend is generally correct.

The stability constant of a complex does not, according to Eq. (1.27), dependon the concentrations of the species comprising the complex. For very dilute solu-tions, however, complexes become less stable than expected from their “literature”stability complex. The reason for this lies in the fact that the equilibrium shown inEq. (1.26) is not strictly correct; a more accurate representation would be

Cd(H2O)62� � 4NH3D Cd(NH3)4

2� � 6H2O (1.30)

(hydrolysis and hydration of ammonia and ammonium ion is ignored, although foran accurate representation, it should be considered—it will not affect the argu-ment). Since ammonia is a much stronger ligand than water (water of hydrationcan be considered as a ligand), it will exchange all the water as long as the am-monia concentration is not too low. If it is very low, then not all the water will nec-essarily be exchanged, and a different equilibrium (or mixture of equilibria) with


a different equilibrium constant will exist. A classic example of this effect is givenby the cobalt complex with thiocyanate, SCN�:

Co(H2O)62� � 4SCN�D Co(SCN)4

2� � 6H2O (1.31)

pink blue

The pink color of the hydrated Co2� ion turns blue when a high concentration ofSCN� is added; if diluted with water, this solution reverts to pink. This color tran-sition is reversible. An aqueous solution of Co2� is pink, while anhydrous Co(II)salts are typically blue (a fact well-known to chemists and even to schoolchildrenwho have experimented with invisible inks). The color change to blue on additionof SCN� to a noncomplexed Co2� solution is caused by dehydration of the Co2�

due to exchange of water with the SCN�, a stronger ligand than water. However,SCN� is not a very much stronger ligand than water to Co2�, and therefore a rel-atively high concentration is required to exchange all the water. At intermediateconcentrations of SCN�, mixed aquo-thiocyanato complexes can be formed,which are pink. The stronger the ligand relative to water, the less the concentra-tion required to exchange the water.

It is worth noting that a statistical effect (different combinations of the var-ious complexants) may result in mixtures of complexants binding more stronglyto a cation than would be expected based on the individual stability constants ofthe complexants [2].

If a compound containing more than one cation is to be deposited, com-plexation could be used to offset the difference in Ksp between the individualmetal compounds. As an example, consider the deposition of (Cd,Hg)S. FromTable 1.1, the value of Ksp for HgS is much lower than for CdS. This means thatunder the conditions of CD, where the sulphide ion is slowly formed, we wouldexpect only HgS to form (until almost all the Hg was used up). Some Cd might beincorporated into the deposit by adsorption, but the deposit should, according toconsiderations of solubility product, be predominantly HgS.

The concentration of Hg in the deposit can be decreased by choosing a com-plex (or mixture of complexes) that complexes Hg more strongly than it does Cd. Inthis case, since Hg forms very strong complexes with many ligands, there is a largechoice. Skyllas-Kazacos et al. deposited films of (Cd,Hg)S using a combination ofammonia and cyanide (the latter is a strong complex for both cations, but more sofor Hg) [3]. In addition, the Hg concentration was much smaller than that of cyanide,while the Cd concentration was larger. This means that there was enough cyanideto complex the Hg but not enough for the Cd. A further factor that may have allowedcodeposition of Cd was the use of the chloride anion, which is a moderately strongcomplex for Hg2� but only a weak one for Cd2�. This combination of factors al-lowed codeposition of the Cd and Hg as sulphides, but the concentrations of Hg inthe films were larger (by a factor of ca. 4) than in the deposition solution. The addi-tion of a large concentration of iodide to the deposition solution would probably


have been even better, since iodide is an extremely weak complexant for Cd and avery strong one for Hg. It would therefore have removed most of the free Hg2� ionswhile only changing the Cd2� concentration a relatively small amount.

This judicious use of a complexant to allow codeposition of two cationswith widely differing values of Ksp is, unfortunately, not always useful. An exam-ple is the deposition of (Cd,Zn)S—a material of interest particularly because of itspotential use in photovoltaic cells (see Chapter 9). The stability constants of Cdand Zn ions are in most cases very similar for any particular complex (although,of course, they vary greatly from one complex to another). This reflects thesimilarity of the chemistry of these two ions (and the difference between them andthe Hg2� ion). Therefore complexation is very limited as a means to control theconcentration of one of these ions relative to the other.

The use of complexation to allow codeposition of alloys is well known inelectroplating. The best-known example is that of brass (Cu/Zn) plating, wherecyanide, which is a stronger complex for Cu than it is for Zn, brings the deposi-tion potentials of the two metals, originally far apart, to almost the same value.There is a direct connection between this effect and the equivalent one for CD.This arises from the fact that, for both CD and electrodeposition of alloys (we in-clude mixed metal compounds in the term alloy), the effect of the complexant isto lower the concentration of free cations. For CD this affects the depositionthrough the solubility product, while for electrodeposition it affects the depositionpotential through the Nernst equation:

E � E0 � �RnF

T� ln �

[[ROexd]]

� (1.32)

where the oxidized species, Ox, the cation in this case, is reduced in concentration,resulting in a more negative deposition potential, E, compared to the standard po-tential, E0. In the case of metal electrodeposition, the reduced species, Red, is themetal that, since it is a solid, can be taken as unity concentration. n is the numberof electrons transferred per molecule of reaction (e.g., for Cd deposition fromCd2�, n � 2) and F is Faraday’s constant (ca. 96,500 coulombs/mole).

The shift in potential due to complexation, �E, (�E � E0) can be approxi-mated by

�E (in mV) � 60 log [cation]/n (1.33)

From Eq. (1.27), we can write

log Ks � log [complexed cation] � log [cation] � a log [ligand] (1.34)

where a is the number of ligand molecules in the complex.Combining Eqs. (1.33) and (1.34) we get

log Ks � �n �E (in mV)/60

� log [complexed cation] � a log [ligand](1.35)


For all but very weak complexes, the concentration of complex ion (and often alsothe free ligand) is normally very much larger than that of the free cation. With thisin mind, Eq. (1.35) can often be approximated by

log Ks � �n �E (in mV)/60 (1.36)

This [or, more accurately, Eq (1.35)] allows us to calculate values of the stabilityconstant of a complexant from tables of electrochemical potentials. For example,a shift of 300 mV in potential due to complexation gives an (approximate) valuefor the stability constant of that complex of 105 (for n � 1) or, for the more com-mon case in CD, where divalent cations (n � 2) predominate, 1010.

1.2 NUCLEATION AND GROWTH

CD can occur either by initial homogeneous nucleation in solution or by het-eronucleation on a substrate, depending on the deposition mechanism (see Chap-ter 3). For this reason, we consider both types of nucleation.

1.2.1 Homogeneous Nucleation

According to simple solubility considerations, a precipitate will be formed whenthe product of the concentrations of anions and cations exceeds the solubilityproduct. From another viewpoint, phase transformation occurs when the free en-ergy of the new phase is lower than that of the initial (metastable) phase. However,there are many examples where the ion product exceeds Ksp, yet no precipitationoccurs—the phenomenon of supersaturation. The solubility product also does notprovide information on how the particles of the precipitate form—nucleation. Nu-cleation involves various physical processes, and both thermodynamic and kineticaspects must be considered.

Homogeneous nucleation can occur due to local fluctuations in the solution—whether in concentration, temperature, or other variables. The first stage in growth iscollision between individual ions or molecules to form embryos (embryos are nucleithat are intrinsically unstable against redissolution—see later). Embryos grow by col-lecting individual species that collide with them. While these species may be ions,atoms, or molecules in general, for CD, adsorption of ions on the embryo seems tobe the most probable growth mechanism. They may also grow by collisions betweenembryos; however, unless the embryo concentration is large, this is less likely.

These embryos may redissolve in the solution before they have a chance togrow into stable particles (nuclei). Because of the high surface areas, and there-fore high surface energies of such small nuclei, they are thermodynamically un-stable against redissolution. They may, however, be kinetically stabilized by lowtemperatures, which increase their lifetime, possibly enough for them to grow toa size where they are thermodynamically stable. This is an important reason why


smaller particles can be formed at lower temperatures in a precipitation reaction;the subcritical embryos last long enough to grow into stable particles, while athigher temperatures they would redissolve, reducing the density of nuclei. This re-sults in an increase of the particle size, since there is more reactant per nucleus.

The critical radius, Rc, is the size where the embryo (nucleus) has a 50:50chance of either redissolving or growing into a stable nucleus; it is determined bythe balance between the surface energy required to form the embryo,

Es � 4R2� (� is the surface energy per unit area) (1.37)

and the energy released when a spherical particle is formed,

Ev � 4R3�L/3

(� is the density of the solid and L is the heat of solution).(1.38)

This balance is shown in Fig. 1.1. The typical size of Rc is about 100 molecules—between 1 and 2 nm in diameter. Solvent molecules can adsorb on the embryosand change their surface energy; the critical radius will therefore depend not onlyon the material of the nucleating phase but also on the solution phase.

1.2.2 Heterogeneous Nucleation

In heterogeneous nucleation, subcritical embryos (or even individual ions) canadsorb onto the substrate. The energy required to form an interface between the

FIG. 1.1 Energetics of nucleation. The critical radius, Rc, depends on the balance be-tween surface and volume energies of the growing particle.


embryo and the solid substrate will usually be less than that required for homoge-neous nucleation, where no such interface exists. Therefore heterogeneous nucle-ation is energetically preferred over homogeneous nucleation and can occur nearequilibrium saturation conditions, compared with the high degree of supersatura-tion often required for homogeneous nucleation. These subcritical nuclei cangrow, either by surface diffusion or by material addition from solution. It shouldalso be noted that nuclei that are subcritical in solution may be supercritical whenadsorbed on a substrate. This is a consequence of reduced contact between nucleusand solution as well as stabilization of the adsorbed nucleus. These processes areshown schematically in Fig. 1.2.

It was noted earlier that even individual ions may adsorb onto a surface.More specifically, depending on the surface chemistry of the substrate, individualions or molecular species may actually be chemisorbed, creating a nucleus for re-action and further growth.

Pure homogeneous nucleation is probably less common that might appearfrom the above discussion. Because of the greater ease of nucleation on a solidphase than homogeneously, any solid matter in the solution will act as a nucleationcenter. It is difficult to prepare solutions without some solid phase (usually dustparticles)—careful filtering is necessary to attain such particle-free solutions. Thatthis is so can be seen from the simple test of shining a laser beam (preferably agreen or blue laser, since scattering is greatly enhanced compared with a red one)through an visibly “clear” solution; the resulting scattering by dust particles is al-most always evident.

1.2.3 Crystal Growth

Once (stable) nuclei have formed, there are several ways in which they can in-crease in size. One is a continuation of the process of embryo growth discussedearlier: adsorption of ionic species from the solution onto the nucleus. Crystal

FIG. 1.2 Processes involved in heterogeneous nucleation on a surface.


growth of this type can be considered a self-assembling process. Thus for CdS, ei-ther Cd2� or S2� will adsorb (as discussed later, since a crystal, and in particulara polar one, is made up of different faces, the adsorption properties of each maybe different, and therefore both types of ions may adsorb on any one crystal). Thenext growth step will then be adsorption of the oppositely charged ion to give anadditional CdS molecule. This process can continue until either all the ions of anyone type are used up or growth is blocked, e.g., by aggregation or by blocking ofthe crystal surface by a foreign adsorbed species. Also, growth may continue butin a different geometric orientation, giving rise to twinning, polycrystallinity, etc.

Another mechanism for crystal growth is known as Ostwald ripening. If asmall nucleus or embryo is close to a larger crystal, the ions formed by (partial)dissolution of the smaller, less stable crystal can be incorporated into the largercrystal. As the smaller crystal becomes even smaller, its dissolution will becomeever more favorable and eventually it will disappear. The result is that the largercrystals grow at the expense of the smaller ones.

If the concentration of particles is sufficiently high, then the probabilityof collisions between these particles becomes high. This can result in eitheraggregation or coalescence. When two particles approach each other, the van derWaals force of attraction (see section 1.3.1) between them will often cause themto stick together. This can continue until a large particle (large in relation to theoriginal particle size) comprising the individual particles has formed (Fig. 1.3A).This is the process of aggregation, and the resulting large particle is called anaggregate. (In colloidal chemistry, the alternative terms of flocculation andfloc are often encountered.) The properties of the aggregate may be similarto those of the individual particles in some ways (such as X-ray diffraction

FIG. 1.3 Aggregation (A) and coalescence (B) of individual particles.


peak broadening, quantum size effects) and very different in others (e.g., lightscattering, sedimentation).

In an aggregate, there are grain boundaries between individual crystallites.However, in some cases, particularly if the temperature is high enough to allowappreciable diffusion of the crystal atoms, surface diffusion may occur where two(or more) particles have aggregated, resulting in the formation of a neck. This istermed coalescence. Coalescence may continue until one large particle is formedfrom the original two or more particles (Fig. 1.3B).

1.2.4 Particle Size Distribution

If nucleation occurs in a very short time, whereas growth occurs separately, oftenover a much longer time but without further nucleation, then the size distributionis likely to be narrow, since all the original nuclei should be of similar size andgrow at the same rate. The opposite case, where nucleation and growth occur si-multaneously, usually results in a wide size distribution.

Homogeneous nucleation normally requires a supersaturated solution, whilegrowth can occur close to the saturation concentration. Therefore rapid nucleationcan occur if supersaturation is rapidly reached. This nucleation lowers the con-centration of reactants below that needed to cause further nucleation. If one of thereactants is supplied at a low concentration after nucleation has occurred (such asby in situ homogeneous formation in the solution), then growth can occur withoutfurther nucleation, resulting in a narrow size distribution.

In CD, where the reaction is slow, it might be expected that nucleation andgrowth will always occur together, resulting in a relatively wide size distribution.This is indeed expected for heterogeneous nucleation on a substrate. However, formechanisms where homogeneous nucleation of an intermediate phase occursrapidly in the solution but conversion to the final compound is a slow process, nu-cleation and growth can still be separated.

1.3 FORCES BETWEEN PARTICLES ANDBETWEEN PARTICLES AND SURFACES

Once nanoparticles have been formed, whether in an early state of growth or in amore or less final size, their fate depends on the forces between the individual par-ticles and between particles and solid surfaces in the solution. While particles ini-tially approach each other by transport in solution due to Brownian motion, con-vection, or sedimentation, when close enough, interparticle forces will determinetheir final state. If the dominant forces are repulsive, the particles will remain sep-arate in colloidal form. If attractive, they will aggregate and eventually precipitate.In addition, they may adsorb onto a solid surface (the substrate or the walls of thevessel in which the reaction is carried out). For CD, both attractive particle–sur-


face and particle–particle forces are required for film formation. In all the stagesof the CD process, except for a very few studies, that of adhesion of the film to thesubstrate is probably the least understood; it is rarely even considered in mecha-nistic studies. Why do some particles stick to a particular substrate and others not?Also (and easier to understand, at least intuitively), why do the particles stick toeach other in building up a film?

In order to understand these sticking phenomena so crucial for the CD pro-cess, we consider the various forces involved—repulsive as well as attractive—involving the particles. We discuss first the more obvious forces and then someless obvious ones that nonetheless may be important in some cases. Since thedominant force in CD is the van der Waals attraction, we will begin with this in-teraction.

1.3.1 van der Waals Forces

The main interaction that determines whether and to what extent particles will ad-here to each other and also (if there is no specific chemical interaction) to a sub-strate is, in most cases, the van der Waals interaction. The van der Waals force ofattraction is a universal interaction that operates between all particles, whetheratoms, molecules, clusters, charged, or noncharged. The attraction is due to an in-duced dipole–induced dipole interaction between particles. The dipoles arise fromfluctuations in electron density around the ion cores, resulting in transient changesin the charge density distribution. This transient dipole in one particle induces anequal and opposite dipole in the other one, resulting in an attraction. It may bethought that all the transient dipoles in the randomly orientated particles wouldcancel each other and average to zero. This is indeed the case for an ensemble ofparticles. However, the correlation between the dipole in one particle and the in-duced dipole in another at any time is not zero—the correlations, and therefore theattractive interactions, do not average out to zero. Such charge fluctutations are auniversal property of matter and occur even in a completely nonpolar material. Ifthere is a permanent dipole in (some of) the interacting particles, these dipoles willalso contribute to the van der Waals interaction. For purely nonpolar particles, theinteraction is known as the London, or dispersion, energy.

The van der Waals interaction between atoms or molecules, E, varies as theinverse sixth power of the distance, d, between them:

E (1.39)

(The minus sign signifies an attractive interaction.) For macroscopic (this includesmicroscopic and nanoscopic) bodies, this interaction is much less short-range, andthe distance dependence varies both with the geometry of the interacting bodiesand with the distance of separation. For macroscopic bodies, it is usually assumed

�1�d6


that the interactions between all the different bodies are additive. Table 1.3 showsvalues of this interaction for various geometries of two interacting bodies relevantto CD. These can be divided into particle–particle attraction (formation andgrowth of aggregates) and particle–plane surface (i.e., the substrate) interaction.Two different distance scales are shown, depending on whether the separation isconsiderably larger or smaller than the radius of the particle. Clearly there willalso be intermediate separations, where the separation and radius are comparable,with intermediate dependence on the separation. Cases where two interactingspheres are identical or of different size are also shown. Both are relevant for CD,where initial aggregation will occur between elementary colloids of approxi-mately the same dimensions but further interaction can occur between two parti-cles of very different sizes.

The Hamaker coefficient, A, is a measure of the interaction and is dependenton the material of the particle as well as on the surrounding medium. Heavy atoms,which are generally more polarizable (i.e., the electron distribution can be more

TABLE 1.3 Interaction Energies and Forces of Attraction Between TwoBodies with Different Geometries

Geometry Energy Force

� ��32

3

A

d

r713r2

3

�

��1

9

6

d

A6

r6

� ��3

3

2

d

A7

r6

�

��6(r

A

1

r

�

1r

r2

2)d� ��

6(r1

A

�

r1r

r2

2)d2�

��1

A

2

r

d� ��

12

A

d

r2�

��2

9

A

d

r3

3

� ��2

3

A

d

r4

3

�

��6

A

d

r� ��

6

A

d

r2�

A is the Hamaker constant, r (r1, r2) the radius of the spherical particles, and d thedistance between surfaces of the two bodies. Note that larger particles will interactmore strongly (more adherent films?).

16Ar31r3

2�

9d6


easily perturbed), generally have a larger value of A than lighter atoms and there-fore a greater attractive interaction. A has units of energy and values that vary typ-ically from several times the thermal energy, kT, to several tens of kT in air or avacuum and typically an order of magnitude less in liquid media.

At large distances between particles, correlation between fluctuations in oneparticle and the induced dipole in another breaks down. This occurs when the timetaken for the interaction (acting at the speed of light) is comparable to the charac-teristic scale of the electron fluctuations, viz. the plasma frequency. The plasmafrequency ranges typically from 10 eV down to 2 eV (closer to the former formany dielectrics and to the latter for metals), which translates into a length scaleof between 600 and 100 nm. At this distance scale, the (at this point, very weak)van der Waals forces are termed retarded forces, because of the appreciable timerequired for the transmitting dipole electromagnetic field to reach the receivingspecies. At this distance, the attraction falls off approximately as the inverse sev-enth power of the distance. It is probable that diffusional and convective motionand electrostatic interactions will dominate at such distances and the van derWaals interaction will be negligible in most cases.

These various relationships between force and particle separation imply thatthe attractive force between particles will become infinite when they touch. In re-ality, other short-range forces will modify this relationship when r is very small,in particular the repulsion from overlap of atomic orbitals. The van der Waals at-traction will then be balanced by this overlap repulsion. At these short distances(a few tenths of a nanometer), the van der Waals attraction will be strong enoughto hold the particles fairly strongly together. This balance between van der Waalsforces of attraction and overlap repulsion forces is shown schematically in Fig.1.4, where the very steep repulsive interaction at atomic distances is due to theoverlap repulsion. Hydration forces (see section 1.3.3) may also result in repulsionbetween surfaces at somewhat greater separations.

Particle adhesion occurs when the distance between bodies is that of anatomic spacing. From Table 1.3, the force between a sphere of radius r and a flatsurface at close approach is

(1.40)

At contact, d is the atomic spacing. For a solid where the van der Waals forcesdominate, the work needed to separate two unit areas from contact to infinity isgiven by

(1.41)

where z is the atomic spacing [� d in Eq. (1.40) at contact]. This is the energy re-quired to produce two new surfaces, i.e., 2�, where � is the surface energy of the

A�12z2

�Ar�6d2


solid. Therefore,

� � (1.42)

Substituting Eq. (1.42) into Eq. (1.40) gives the force of adhesion at contact interms of the surface energy:

F � 4r� (1.43)

This relation is clearly very simplified, being based on a number of approxima-tions, such as the validity of the use of the Hamaker constant at such close dis-tances and the particle and surface being of the same material. Also, the relation-ship between surface force and van der Waals forces does not hold for manysolids, in particular for metals where metallic bonding is important. Nonetheless,if taken as an indication of the forces holding particles to each other and to sur-faces, it does give a feel for these forces.

1.3.2 Electrostatic Forces

For solid particles dispersed in a liquid medium, there exists, in most cases, a layerof charge separation at the phase boundary—the electrical double layer. A num-

A�24z2

FIG. 1.4 Resultant interaction energy between two particles with van der Waals attrac-tive interactions and electron overlap repulsion interactions.


ber of processes can cause this double layer. One of the most common is adsorp-tion of charged species at the solid–liquid interface. For example, in a colloidal solof CdS prepared by precipitation of a Cd salt with sulphide ions, Cd2� or S2�

(HS�), depending on which is in excess, will adsorb at the CdS surface. Anothermechanism for formation of surface-charged species is surface dissociation. Acommon example of this is the case of metal oxides in water; the water may dis-sociate at the oxide surface as follows:

(1.44)

leaving a positively or negatively charged surface. Such a reaction will, obviously,be very pH dependent.

Yet another possibility for formation (or change) in a double layer is by ac-cumulation of one charge type. This may occur by doping the solid with an ion ofvalence different from that of the solid (e.g., In3� in CdS) or by illumination withsuper-bandgap illumination. In the first case, the (in this case) electron donor(In3�) is immobile, while the donated electron is mobile; if the electron is trans-ferred to the liquid, then the solid will become positively charged. The same oc-curs for an illuminated solid where electron/hole formation occurs, if one of thecharges is preferentially injected into the liquid. Even in the absence of charge in-jection into the liquid, localized (e.g., trapped) charges will affect the double layerif the countercharge is (relatively) delocalized over the particle. The double layerclose to a (near) surface-localized donor will be different than that for the rest ofthe particle. Such an effect is probably not important in large particles, where suchfluctuations can be evened out. For very small particles, however, where only asingle “dopant” may exist, this effect may be appreciable. Even for a pure and per-fectly stoichiometric particle, the double layer need not be homogeneous aroundthe particle. For example, a CdS particle will consist of different crystal faces.Most notably, the opposite polar faces, consisting of only Cd or of S atoms, canbe expected to possess different double layers. Incidentally, these polar facesmight be expected to attract other polar faces of opposite polarity and repel thoseof the same polarity. Such an effect would lead to some form of self-assembly.However, any effect of this nature will be much smaller (if it exists at all) in solu-tion compared to vacuum or air, due to adsorption of ions from solution onto thepolar faces, which will tend to neutralize this effect.

The charge at the surface of a solid (including any adsorbed species) will bebalanced by a countercharge in the electrolyte; the double layer as a whole is elec-trically neutral. The countercharges remain in the vicinity of the surface-adsorbedcharge but, due to thermal motion, do not accumulate at the surface but move in amore or less diffuse cloud surrounding the particle. The extent to which this layer

MMO � H2OD

D

MOH� � OH�

MOMOH� � H�


of countercharge extends into the solution—how diffuse it is—depends on theconcentration of charged species in the solution. This gives the solution screeninglength. The screening length is analogous to the space charge layer width in asemiconductor; in the same way, the screening length, �, is a function of thesquare root of the charge (ionic) concentration.

� � � �1/2

(1.45)

where z is the ionic charge, n is the ionic concentration, �0 is the permittivity offree space, and � is the dielectric constant of the material. (The dielectric constantis normally taken as constant. It should be pointed out, however, that for nanopar-ticles of several nanometers or less, the value of � decreases with particle size, as-suming the particles are in a medium of smaller dielectric constant than the parti-cles themselves, a reasonable assumption for our purposes. This effect is treatedin some detail by Lannoo et al. [4].) The thickness of this diffuse layer (alsoknown as the Gouy layer) is the inverse of �; i.e., the potential drop across the dif-fuse layer, �d, decays to �d/e (where e in this case is the natural exponent, 2.718)over a distance ��1.

Figure 1.5 shows a schematic representation of the double layer at a planarsolid–liquid interface. The potential drop across the Helmholz layer is shown aslinear (in the presence of specific adsorption, it will not be completely linear), fol-lowed by a tailing-off of the potential into the diffuse layer. For concentrated so-lutions (�0.1 M) the diffuse layer is typically a nanometer or less, while for dilutesolutions it may be tens or even hundreds of nanometers.

2e2z2n��0kT

FIG. 1.5 Schematic diagram of the electrical double layer at a solid–liquid interface.


The crucial importance of the double layer when dealing with colloidal par-ticles dispersed in a solution is due to the repulsion of one particle by another.While overall the particles are neutral, because the diffuse layer can extend intothe solution, the unbalanced charge in the diffuse layer of one particle experiencesa repulsion by that of another particle. Normally, from the Coulomb law of elec-trostatics, the force between two equal (in charge and in sign) particles is given by

F � (1.46)

where F is the force acting between two charges, q1 and q2, separated by a distanced. Because of the presence of the diffuse layer, however, the repulsion force be-tween two particles is strongly dependent on the screening length, �, and is ap-proximately proportional to exp (-�d); the force of repulsion between two colloidswill decrease exponentially with distance.

Overlap of the diffuse layers of approaching particles prevents them fromgetting close to each other, and the particles form a stable colloidal solution. Morerelevant for our purposes, for a moderately concentrated electrolyte of the typenormally encountered in CD (on the order of 0.1 M or more total concentration),this diffuse layer is around 1 nm or less. The diffuse layer screens the surface chargeand allows the like-charged particles to approach each other closely, to the pointwhere the van der Waals forces of attraction dominate, causing aggregation. This is the basis of salting-out of a colloid; addition of a strong electrolyte to the colloidreduces the thickness of the diffuse layer, allowing closer approach of the particlesto each other and eventual aggregation and precipitation. The competition betweenthe attractive van der Waals and repulsive electrostatic forces and the importance ofthis competition in colloid stability is known as the DLVO theory, named for thescientists who developed a theoretical analysis of the overall interaction (Derjaguin,Landau, Verwey, and Overbeek). The resultant interaction is shown in Fig. 1.6. Fora dilute electrolyte, there is a relatively large barrier to aggregation where thedouble-layer repulsion dominates the interaction to the greatest extent and theinteracting particles fall into what is known as the secondary minimum. At thispoint they are kinetically stabilized against aggregation. As the electrolyte concen-tration increases, the barrier becomes smaller and eventually disappears, resultingin the particles becoming trapped in the primary minimum, i.e., aggregation.

It is worth noting that the electrostatic force can be attractive as well as re-pulsive, depending on the sign of the two charges. For the case where a single col-loidal species is present in the solution, it will be repulsive, since all the particleswill have the same charge. Two different colloidal species of opposite chargecould conceivably be present in the deposition solution, either because the CDprocess involves conversion of one species into another (e.g., metal hydroxide andmetal chalcogenide) or because two or more different cations, or even differentvalence states of the same cation, are present.

q1q2�4�0�d


1.3.3 Entropic and Other Short-Range Forces

Apart from double-layer, van der Waals, and electron-overlap interactions, anytwo bodies in a liquid medium (even if only one, or even neither, is charged) willexperience a (usually) repulsive component of force as they approach each other.This is due essentially to entropic considerations. As two bodies approach eachother very closely, the species in solution have increasingly less room in which tomove; the entropy of these species therefore decreases, producing a repulsiveforce between the two bodies. In its simplest form, this force is usually consideredto exist between two infinite flat plates. In the context of forces between colloidalparticles of the type common in CD, this entropic force should be less important,since the solution species between two approaching particles can be relatively eas-ily pushed out of the intervening space, both due to the small size and due to theapproximately spherical shape common in these systems. In fact, in this case theremay even be a weak attraction—the depletion interaction—due to the smaller den-sity of solution species in the space where the two particles are closest to eachother (from which they have been pushed out) and the surrounding solution. Theentropic force of repulsion may be important, however, when considering thesticking behavior of larger aggregates at plane surfaces.

FIG. 1.6 DLVO interactions showing the energetics of colloidal particles as a competi-tion between electrostatic double-layer repulsion and van der Waals attractions. The pri-mary minimum is due to strong short-range electron overlap repulsion (shown in Figure 1.4but not shown here).


This entropic force is important where adsorption of polymers occurs oncolloidal particles. This is due to interaction between polymer chains on the inter-acting particles: As the particles approach each other to the point where the poly-mer chains of the two particles interact, there is an decrease in entropy due to con-finement of the chains, in an analogous manner to the solution species discussedearlier, with the same result—repulsion. This is the basis of polymeric stabiliza-tion of colloids; it is generally undesirable in CD, since adhesion and aggregationare preferred in this case. However, in view of the fact that the presence of suchpolymers (and other stabilizing adsorbates) may prevent the aggregation neededto build up a CD layer, it is important to be aware of the effect.

There are other close-range forces related to entropy changes, including var-ious interactions between solution species and a solid surface, such as solvation(in water, hydration) forces. Hydration forces can occur when hydrated cations areadsorbed at interacting surfaces. As these surfaces approach each other closely,loss of water of hydration is necessary in order to allow closer approach. Whilethese forces can be repulsive, attractive or oscillating, they are most likely to berepulsive under the conditions of CD. Such forces may be very important for CD,which is almost always carried out in the presence of a high ionic concentration.For example they could be a cause of poor adhesion of some CD films. Solvationforces are treated in detail in Israelachvili’s book—see Further Reading at the endof this chapter, “Forces” subsection.

If this treatment of forces between particles fails to convince the reader thatit is natural for particles to stick together, one can resort to the more intuitive ap-proach. It is well known that inorganic colloids require a stabilizing agent to pre-vent their sticking together and eventually precipitating. In other words, precau-tions usually have to be taken to prevent the natural tendency of these particleseventually to stick to each other.

1.4 CHARACTERIZATION TECHNIQUES—SOME CAUTIONS

Many techniques have been used to characterize CD films. The purpose of thissection is not to review all these techniques, but only to draw attention to some ofthem that are sometimes misinterpreted.

Three common techniques used are transmission electron microscopy to-gether with electron diffraction, powder X-ray diffraction, and optical absorption(or transmission) spectroscopy.

1.4.1 Transmission Electron Microscopy/ElectronDiffraction and X-Ray Powder Diffraction

Transmission electron microscopy (TEM) is used to image nanocrystal (lateral)size, shape and size distribution. Electron diffraction (ED) provides information


on the composition of the deposit, crystal phase, and orientation. X-ray diffraction(XRD) also provides similar information on composition (more accurately thanED) and phase as well as crystallite orientation. In the last, ED is superior in manyways, since a much smaller area can be selected (selected area diffraction—SAD)and, in addition, azimuthal lattice alignment between deposit and substrate (epi-taxy) can be determined from ED but not from the commonly used powder XRDmeasurements; powder XRD reveals texturing (one particular crystal face parallelto the substrate for all crystals for perfect texturing) but not orientation (crystal lat-tices in any direction parallel to the substrate of all crystals aligned in the sameway). There are XRD measurements that can distinguish orientation, but while be-coming somewhat more common, these are still rather infrequently used, at leastin CD studies, compared to the normal “�–2�” measurement.

CD films are often nanocrystalline. One very important use of XRD whendealing with nanocrystals is to estimate crystal dimensions through the Scherrerrelationship:

crystal diameter � (1.47)

where � is the X-ray wavelength (0.1541 nm for Cu K� radiation, a commonlyused source), �(2�) is the peak full width at half maximum (FWHM) in radians,and � is the peak position.* The shape of the crystal can also modify this relation-ship, which is valid for a spherical crystal (close to the shape often encountered).

As a rough and useful rule of thumb, a peak FWHM of 1° at an angle of 2� � 25° (a common approximate position) means a crystal size of ca. 12 nm (forCu K� radiation), and the size is inversely proportional to the FWHM. Actually, tobe more precise, what is measured is not necessarily crystal size but coherencelength, the length over which the periodicity of the crystal is complete. An exam-ple of a coherence length smaller than the crystal size is a twinned crystal; XRDmeasures the size of each individual twin. Other causes for XRD peaks beingbroader than expected based on crystal size is the presence of strain in the crystalsor other defects, such as dislocations, which destroy the long-range lattice order.Separation between crystal size and strain can be made if several different peaksare present, since the angular dependences of the two factors are different (see Ref.5 for an example of this). Thus, the interpretation of XRD peak broadening shouldbe carried out with care and preferably using complementary TEM measurements.The opposite case, where the peaks are narrower than expected based on crystalsize, does not occur; a narrow peak means a (relatively) large coherence length and

1.3��(2�) cos �

* Note that in the common �/2� measurement, the � in cos � is half of the 2� value. For example, if thepeak being measured is at 2� � 25°, cos � will be cos 12.5°. For small angles, the error in taking cos2� instead of cos � is not too large (�10% at 2� � 25°, for example). However this error becomeslarger as the peak angle increases.


therefore crystal size. However, even here, interpretation is not always straightfor-ward. For example, the XRD pattern of a deposit of tall cylinders of small crosssection will give a peak width characteristic of the height but not of the cross sec-tion (the latter will be seen in TEM images). Thus the TEM and XRD sizes willnecessarily be different in such a case. Another example is where there is a mixtureof large and small crystals. Even if the large crystals constitute a relatively smallfraction of the total material, they may in some cases dominate the XRD pattern,since peak heights decrease with decreasing crystal size due to increase of peakwidth and (ideally) constant peak area for the same quantity of material.

If the crystal size becomes very small (a few nanometers), the XRD peakswill be very wide and also relatively weak. There is no shortage of examples in theliterature where samples have been classified as “amorphous” or “poorly crys-talline” either on the basis of the lack of any XRD peak or because the peak(s)were very broad. When carrying out an XRD measurement on a CD film, in par-ticular, a particularly thin one (some tens of nanometers or less), if no peaks areseen in the measurement, it is advisable to repeat the measurement over a narrowrange (where a major peak is expected) and with a very slow scan (e.g., 10°/hr oreven slower). If a thin-film attachment is available, this will reduce the likelihoodof such misinterpretations. It is useful to remember that except for compounds thatare commonly amorphous, CD semiconductor films are rarely truly amorphous.

1.4.2 Optical Absorption

Optical absorption spectroscopy is often carried out on CD films to verify that thefilms have a bandgap expected from the deposited semiconductor. Additionally,since CD films are often nanocrystalline and the most apparent effect of verysmall crystal size is the increasing bandgap due to size quantization (the effect isvisible to the eye if the bandgap is in the visible region of the spectrum), absorp-tion (or transmission) optical spectroscopy is clearly a fast and simple pointer tocrystal size, since bandgap–size correlations have been made for a number ofsemiconductor colloids and films.

There are some potential problems that should be taken into account wheninterpreting such spectra. A spectrophotometer measures transmission (and maybealso reflection) but not absorption. What is measured as absorption is a transmis-sion measurement that is mathematically manipulated to convert it to absorption.Absorption is usually measured as absorbance, A, which by definition is given by

A � log10 Io/I (1.48)

where I is the intensity of the transmitted light and I0 that of the incident light.The transmission, T, is

T � �II0� (1.49)


The spectrophotometer measures the transmission and, if an absorption measure-ment is carried out, converts the transmission into absorbance using these equa-tions. This conversion works fine for samples where there is no reflection, eitherspecular or diffuse, as is the case for nonturbid solutions. However, for films thereis invariably some reflection, which is often quite large, particularly for films ofhigh dielectric constant (or refractive index) materials, such as PbS and PbSe. Ad-ditionally, if the films are not completely transparent, then scattering introducesan extra element of reflection. Therefore, to measure the real absorption of a film,a reflection measurement must also be carried out and correction for this reflec-tion made. The correction will be approximate and depends on the nature of thefilm itself. However, that most commonly used is

Tcorr � �1 �

TR

� (1.50)

where Tcorr is the corrected transmission, T is the measured transmission, and R isthe reflectance. This correction neglects reflection from the film/substrate inter-face (assuming the front face of the film faces the illumination source), and it canbe calculated that this will give a value of Tcorr that is too small (therefore an ab-sorption that is too high). Use of (1 � R)2 in the denominator of Eq. (1.50) in-cludes this reflection but tends to give a value of Tcorr that is too high. For verythin films, where reflection cannot be assumed to originate from a surface (i.e.,where the film thickness is not much greater than the depth from which reflectioncan occur), the calculation is more complicated. Fortunately, reflection from suchthin films is also normally low, and therefore the correction is less important.

To collect scattered transmission and correct for diffuse reflectance, a spec-trophotometer with an integrating sphere should be used. This is important if filmsare not very transparent.

In many cases, the lack of correction for reflection will not affect the shape ofthe optical spectrum very much, but it will just give an inflated value for absorption.However, there are frequently cases where the shape of the absorption spectrum isalso appreciably changed after the reflection correction is carried out. Also, if theprimary absorption is a weak one, then correction for reflection, and in particular forscattering, is crucial, since the absorption may be masked by these effects.

The absorption coefficient, �, of the semiconductor can be derived from theabsorption (or transmission) spectrum according to the Beer–Lambert equationapplied to solids:

I � I0e��t (1.51)

where t is the film thickness.For bulk semiconductors, the relationship between � derived from the ab-

sorption spectrum and the semiconductor bandgap, Eg, is given by:

(�h�)n � C(h� � Eg) (1.52)


where n is 2 for a direct transition and 0.5 for an indirect one and C is a constant.A plot of (�h�)n vs. h� should then give a straight line (over much of the absorp-tion onset region), which extrapolates at a zero value of (�h�)n to the value for Eg.This calculation is based on the density of states in the valence and conductionbands of the bulk semiconductor. For semiconductors in the quantum size regime,however, the density of states may be quite different than in the bulk. Addition-ally, a large size distribution, meaning distribution of bandgaps, will smear out theonset. In practice, however, this extrapolation often appears to give a reasonablevalue of the bandgap. The important thing is to be aware of the limitations of themeasurement. Just because a few points on a plot based on Eq. (1.52) give astraight line does not automatically mean that the bandgap can be obtained fromthe extrapolation of this line; not too infrequently, a weaker absorption onset atlonger wavelengths has been ignored, although consideration of the entire spec-trum, together with the expected behavior of the material, would lead one to con-clude that this weaker absorption determines the bandgap.

FURTHER READING

Solution Chemistry

Just about any textbook on inorganic chemistry. A particularly useful one, whichthe author referred to on many occasions while writing this chapter, is:

G Wulfsberg. Principles of Descriptive Inorganic Chemistry. Mill Valley, CA: UniversityScience Books, 1991.

For more extensive tables of solubility product than given here, Lange’s Hand-book of Chemistry (J.A. Dean, 11th ed. New York: McGraw-Hill, 1999) gives anextensive list. The standard CRC Handbook of Chemistry and Physics (CRCPress) also gives a useful, if less extensive, list. Reference 1 also provides an ex-tensive list of sulphides.

Nucleation and GrowthHC Freyhardt, ed. Crystals: Growth, Properties, and Applications. New York: Springer-

Verlag, 1983, Vol. 9 (this is more mathematical than the other treatments).HK Henisch. Crystal Growth in Gels. University Park: Pennsylvania State University

Press, 1973.AE Nielsen. Kinetics of Precipitation. New York: Pergamon Press, 1964.BN Roy. Crystal Growth from Melts. New York: Wiley, 1992.

ForcesF Evans, H Wennerström. The Colloidal Domain: Where Physics, Chemistry, Biology, and

Technology Meet. New York: VCH, 1994.


PC Hiemenz. Principles of colloid and surface chemistry. New York: Marcel Dekker, 1986.J Israelachvili. Intermolecular and Surface Forces. Orlando, FL: Academic Press, 1992.H Ohshima, K Furusawa, eds. Electrical Phenomena at Interfaces: Fundamentals, Mea-

surements, and Applications. New York: Marcel Dekker, 1998.

REFERENCES1. S Licht. J. Electrochem. Soc. 135:2971, 1988.2. P O’Brien, DJ Otway, D Smith-Boyle. Thin Solid Films 361:17, 2000.3. M Skyllas-Kazacos, JF McCann, R Arruzza. Appl. Surf. Sci. 22/23:1091, 1985.4. M Lannoo, C Delerue, G Allan. Phys. Rev. Lett. 74:3415, 1995.5. SB Qadri, JP Yang, EF Skelton, BR Ratna. Appl. Phys. Lett. 70:1020, 1997.


2General Review

2.1 A (BRIEF) HISTORY OF CHEMICAL DEPOSITION

Chemical deposition (CD) of films is not a new technique. As early as 1835,Liebig reported the first deposition of silver—the silver mirror deposition—usinga chemical solution technique [1]. The first reported CD of a compound semicon-ductor film appears to be formation of “lüsterfarben” (lustrous colors) on variousmetals from thiosulphate solutions of lead acetate, copper sulphate, and antimonytartrate, giving films of PbS, Cu-S or Sb-S, which possessed “splendid” colors (in-terference colors resulting from various thicknesses of the deposited films) [2].More “recent” studies of this general process have invoked an electrochemicalmechanism for some thiosulphate depositions, based on the dependence of depo-sition on either the nature (standard electrochemical potential) of the metal sub-strate or on a contacting non-noble metal (which can be looked at as an internalelectrochemical deposition) [3–5]. However, while it is probable that an electro-chemical or mixed electrochemical/chemical mechanism may be applicable onsome metal substrates, some of these depositions do appear to be true CD pro-cesses. PbS is probably the clearest of these; others were Cu2S, Ag-S, Bi-S, Sb-S.Fe, Ni, and Co all formed apparent sulphide films on Fe substrates, while ammo-nium molybdate deposited a film from a thiosulphate solution that did not containS and was probably an oxide. Beutel and Kutzelnigg cover a wide range of depo-


sitions from thiosulphate solutions—both CD and electrochemical [5]. Only in afew cases were these films characterized other than by color.

In 1884, Emerson-Reynolds reported deposition of PbS films by reactionbetween thiourea (thiocarbamide) and alkaline lead tartrate, where “the metallicsulphide . . . became firmly attached as a specular layer to the sides of the vessel”[6]. A wide range of substrates, apart from that just mentioned (a glass beaker),was successfully used for this deposition; porcelain, ebonite, iron, steel, and brasswere specifically mentioned. Even more important, the deposits were very adher-ent, as quantified by their ability to “withstand considerable friction with a wash-leather, and under this treatment take a fine polish.”

Infrared photoconductivity in CD PbS films was reported nearly a centuryago [7,8], and this application has been a central driving force for subsequentinvestigations in CD lead chalcogenide films. The early literature invariablymentions the pioneering work of Kutscher in Germany, during World War II, indeveloping CD PbS and PbSe films for infrared detectors. However, the apparentlack, in all these references to Kutscher’s studies, of any published papers mightsuggest (to the overly suspicious reader) a possible military involvement in thesestudies. These (and subsequent) studies succeeded to the extent that CD was,and apparently still is, the main technique used in making commercial PbS andPbSe infrared detectors (vacuum evaporation was the only competing technology)[9,10].

For a long time, CD was then essentially limited to PbS and PbSe. It was notuntil 1961 that deposition of CdS, now the most widely studied material in CD, was explicitly reported [11] (although CdS deposited from a thiosulphate so-lution which “sticks obstinately to the glass” was already noted in 1912 [11a]).The range of materials deposited by CD was gradually extended, particularly inthe 1980s, to include sulphides and selenides of many metals, some oxides, andalso many ternary compounds (Tables 2.1 and 2.3 in this chapter list films de-posited by CD).

Chemical deposition received a major impetus after CdS films, chemicallydeposited onto CdTe (and, later, onto CuInSe2) films, were shown to give supe-rior photovoltaic (PV) cells compared with the previously evaporated CdS. Thefirst reference to CD CdS used in thin-film PV cells appears to be from Uda et al.[12], although no special importance was attached to the CD technique in that pa-per. Birkmire et al. showed that CD CdS was as good as evaporated (Cd,Zn)S asthe heterojunction partner in CuInSe2-based thin-film cells, giving 10.6% effi-ciency [13]. Two years later, the efficiency of CuInSe2 cells using CD CdS had in-creased to 12.8% [14]. In 1991, Chu et al. used CD CdS to make high-efficiency(13.4%) CdTe/CdS thin-film cells, explicitly stressing the beneficial role of theCD CdS [15] and followed this a year later with a �14.5% cell using the CD CdS[16]. Nowadays, Cd is almost universally used to form the CdS layer on bothCdTe and CuIn(Ga)Se2 thin-film PV cells.


Another cause of interest in this technique is due to the fact that the crystalsin most as-deposited CD films are very small. Considering the current interest innanoparticles, CD is an excellent technique to deposit nanocrystalline films. Morespecifically, if the nanocrystals are small enough, they exhibit size quantization,the most obvious manifestation of which is an increase in the optical bandgap withdecrease in crystal size, as was shown for CD CdSe [17] and later for CD PbSe[18,19]. In fact, the changes in optical spectra that occurred in these films as afunction of nanocrystal size were exploited to provide information on the differ-ent mechanisms of the deposition process [20].

Chemical deposition has also been emphasized as a technique to form solarcontrol coatings. Solar control coatings are envisaged for use on windows in hot cli-mates and possess the (ideal) characteristic of moderate to high visible transmissionto provide adequate lighting, together with high infrared (0.7–2.5 �m) reflectanceto minimize heating by solar energy. CD is a potentially suitable method to preparethese coatings on the large areas of glass that would be needed. Most of the work inthis field has been carried out by Nair and Nair in Mexico using various semicon-ductor films, mainly PbS [21–23] and CuxS [23]. See also this group’srecent review on this work [24]. These coatings are normally yellowish or neutralby transmitted light and various shades of gold, blue, or purple by reflected light.

2.2 WHAT IS CHEMICAL DEPOSITION?

Chemical deposition refers to the deposition of films on a solid substrate from areaction occurring in a solution (almost always aqueous). Using the prototypicalCdS as an example, a Cd salt in solution can be converted to CdS by adding sul-phide ions (e.g., as H2S or Na2S); CdS immediately precipitates (unless the solu-tion is very dilute—a few millimolar or less, in which case CdS often forms as acolloidal sol). Another pathway for CdS formation, one that does not require freesulphide ions, is decomposition of a Cd-thiocomplex (a compound that binds toCd through a sulphur atom). In CD, the trick (or at least one of them) is to controlthe rate of these reactions so that they occur slowly enough to allow the CdS ei-ther to form gradually on the substrate or to diffuse there and adhere either to thesubstrate itself (at the early stages of deposition) or to the growing film, rather thanaggregate into larger particles in solution and precipitate out.

This rate control can be accomplished by generating the sulphide slowly inthe deposition solution. The rate of generation of sulphide, and therefore reactionrate, can be controlled through a number of parameters, in particular the concen-tration of sulphide-forming precursor, solution temperature, and pH.

The CdS forms through a number of different possible pathways: simpleionic reaction between Cd2� and sulphide ion; topotactic conversion of Cd(OH)2,which may be present in the deposition solution, to CdS by sulphide; and decom-position of a complex between Cd (whether as a free ion or as a Cd compound,


e.g., Cd(OH)2) and the sulphide precursor (often thiourea, which, like otherchalcogenide precursors, also acts as a complexant for metal ions).

Although CD can be carried out in both acidic and alkaline solutions, mostCD reactions have been carried out in alkaline solutions. This is necessary for se-lenide deposition using selenosulphate (see later), which is unstable in acid solu-tion (of course, the chalcogenide precursor must not be too stable under all condi-tions, otherwise they will not work). Therefore to prevent (at least bulk)precipitation of metal hydroxides in the deposition solution, the metal ion must becomplexed. There is a very wide range of possible complexing agents available;the most used are intermediate in complexing strength—not too weak, in order toprevent bulk precipitation of hydroxide, but not too strong, which may prevent de-position of the desired film altogether.

2.3 WHAT MATERIALS CAN BE DEPOSITEDBY CD?

In principle, CD can be used to deposit any compound that satisfies four basic re-quirements.

The compound can be made by simple precipitation. This generally, al-though not exclusively, refers to the formation of a stoichiometric com-pound formed by ionic reaction.

The compound should be relatively (and preferably highly) insoluble in thesolution used (except in a very few cases, this has been water).

The compound should be chemically stable in the solution.If the reaction proceeds via the free anion, then this anion should be rela-

tively slowly generated (to prevent sudden precipitation). If the reactionis of the complex-decomposition type, then decomposition of the metalcomplex should similarly occur relatively slowly (see Sec. 2.5 for a de-scription of reaction mechanisms).

Of course there are other specific factors that need to be taken into account, par-ticularly whether the compound will form an adherent film on the substrate or not.However, the preceding four factors are general requirements.

2.4 FORMATION OF THE (CHALCOGENIDE)ANIONS

One of the requirements for CD is either slow release of the anion—in most casesa chalcogenide anion—or slow decomposition of a suitable complex containing achalcogenide atom. This section is confined to the former and discusses slow anionrelease. The precursors used up to now, together with some of the reactions leadingto anion formation, will be briefly described (more details can be found in Chap. 3).


2.4.1 Oxide

Many oxides have been deposited by CD using many different techniques, someof which are described below. Section 2.9.4.1 provides a wider overview of thesemethods. Often, what is deposited is a hydroxide or hydrated oxide. In manycases, hydroxide ion is generated slowly. There are a number of methods to dothis, the most common being hydrolysis of urea:

(NH2)2CBO � 2H2O → (NH4)2CO3 (2.1)

where the carbonate formed partially hydrolyses to give OH�. Another techniqueis reduction of nitrate to nitrite by an alkylamineborane:

NO3� � H2O � 2e� (from alkylamineborane) → NO2

� � 2OH� (2.2)

Additionally, hydroxide may initially be present, but the reaction may be sloweddown e.g. by complexation of the metal ions or by carrying out the reaction at lowtemperature. The hydroxide reacts with the metal ion to give the metal hydroxide,a hydrated oxide or an oxide, depending on the chemistry of the particular metal-(hydr)oxy system. The hydroxides or hydrated oxides can be heated in air or oxy-gen to form the oxides.

Another method used to deposit oxides, particularly those with a higheroxidation state than the starting cation, uses persulphate, S2O8

2�, a strong oxidiz-ing agent. While the exact mechanism of oxide formation using persulphateis unclear, it appears to involve internal electrochemical reactions; e.g., forPbO2:

Pb2� � 2H2O → PbO2 � 4H� � 2e� (2.3a)

S2O82� � 2e� → 2SO4

2� (2.3b)

However, it is possible that free radicals, such as ·OH, are involved, since persul-phate hydrolysis can proceed with formation of H2O2, which is itself sometimesused to deposit oxides:

S2O82� � 2H2O → 2SO4

2� � H2O2 � 2H� (2.4)

2.4.2 Sulphide

Thiourea (SC(NH2)2), the sulphur analogue of urea, is the most commonly usedsulphur precursor. There are a number of possible decomposition routes forthiourea in aqueous solution (it is invariably used in alkaline solutions). Probablythe most important is

SC(NH2)2 � OH�D HS� � CN2H2 � H2O (2.5)

which generates sulphide ion (the cyanamide, CN2H2, can hydrolyze further, butthis need not concern us at present). Actually, aqueous solutions of thiourea are


not very unstable; the presence of a cation that can precipitate an insoluble sul-phide is necessary for the decomposition to proceed at a reasonable rate. This im-plies that reaction (2.5) is actually an equilibrium; the metal ion removes sulphideion, which drives the equilibrium continually to the right.

Thioacetamide (H3C.C(S)NH2) has also been commonly used in CD. It hasthe advantage, compared with thiourea, that it works in both acid and alkaline so-lution. A general decomposition reaction for sulphide formation is

H3C.C(S)NH2 � 2H2O → CH3COOH � H2S � NH3 (2.6)

In alkaline solution, the sulphide will be in the form of sulphide (HS� and S2�)ions.

Thiosulphate (S2O32�) was the original sulphur source in early CD pro-

cesses, and, while less commonly used nowadays, it still has a place in the mod-ern CD literature. It is most commonly used in somewhat acidic solutions, al-though it has also been employed in alkaline solution. Thiosulphate is unstable infairly acidic solutions and decomposes to give elemental sulphur, e.g.,

S2O32� � H� → S � HSO3

� (2.7)

It is often suggested that the thiosulphate, a mild reducing agent, reduces this sul-phur to sulphide.

There are a number of other potential sulphide-forming reactions, depend-ing on pH (see Sec. 3.2.1).

It should be pointed out here that the SMS bond in thiosulphate is easilybroken. In view of the strong complexes it forms with some metal cations, theprobability of a mechanism whereby the SMS bond of the complex is broken,leading to metal sulphide formation without formation of sulphide ion, should beseriously considered (see Sec. 2.5).

2.4.3 Selenide

Selenourea, the selenium analogue of thiourea, which hydrolyzes in the same wayto give selenide ions, was once the most common source of Se. It is an unstablecompound that requires the presence of a reducing agent—usually Na2SO3.Dimethylselenourea is more stable than selenourea but still difficult to work with.The most common Se precursor used nowadays is sodium selenosulphate(Na2SeSO3), which can be considered as the analogue of thiosulphate, with one Satom substituted by Se. It is much more stable (and cheaper) than selenourea andtherefore simplified greatly the deposition of selenides. It can only be used in al-kaline solutions (it decomposes at pH values lower than ca. 7 to precipitate ele-mental red Se). Its alkaline hydrolysis is usually given as

SeSO32� � OH� → HSe� � SO4

2� (2.8)


although it is very probable that SO42� is the end product of a more complicated

reaction. As for most other mechanisms discussed earlier, Eq. (2.8) is probablyalso an equilibrium between selenide and some sulphur–oxygen intermediate,with the metal cation removing the selenide as for the thiourea decomposition inthe presence of metal cations.

2.4.4 Telluride

There are only a very few reports of telluride deposition by CD. This is due toa number of factors: the relative instability of the tellurium analogues of the S-and Se-forming precursors; the strong reducing conditions necessary to form tel-luride ions; and the rapid reaction of telluride ions with oxygen dissolved in thesolution. CdTe films have been formed by using hydrazine—a strong reducingagent—in a deposition solution containing TeO2, and very thin films have beenobserved to form when H2Te was added to a solution of a Cd salt. Recently, thedeposition of CdTe has been reported using the tellurium analogue of seleno-sulphate (tellurosulphate). While Te is apparently only very slightly soluble insulphite, it is apparently enough to deposit tellurides. For both methods, stoi-chiometric films are more difficult to obtain than for sulphides or selenides, withelemental Te typically also formed. A very recent study has also described con-version of CD Cd(OH)2 films to CdTe using a solution of Te in hydroxymethanesulphinic acid.

2.4.5 Halides

Halides (confined at present to silver halides) can be deposited by hydrolyzing awater-soluble halogeno-alcohol (halohydrin) to slowly form halide ions in thepresence of Ag� ions:

X(CH2)nOH � H2OD X� � H� � HO(CH2)nOH (2.9)

The solubility products of most halides are much higher in general than those ofchalcogenides. Those of the silver halides are fairly low, which allows these de-positions to take place readily.

2.4.6 Other Anions

Although CD seems to have been limited to chalcogenides (including oxides andhydroxides) and isolated cases of carbonates, silver halides, and elemental Se, itshould be possible to deposit salts of other anions. There are a number of other an-ions that can be slowly and homogeneously generated. These are discussed inChapter 3.


2.5 MECHANISMS OF CHEMICAL DEPOSITION

2.5.1 Introduction

The mechanisms of CD processes can be divided into two different processes: for-mation of the required compound by ionic reactions involving free anions, and de-composition of metal complexes. These two categories can be further divided intwo: formation of isolated single molecules that cluster and eventually form acrystal or particle, and mediation of a solid phase, usually the metal hydroxide. Weconsider first the pathways involving free anions and defer to later those where ametal complex decomposes.

A starting point for discussing the mechanisms of CD is to consider a sim-ple precipitation reaction. If H2S is added to an aqueous solution of a Cd salt, yel-low CdS precipitates out immediately. H2S precipitates the sulphides of mostcations (the alkaline and alkaline earth sulphides are soluble in water); this is thebasis of the well-known (at least, in the author’s university days) inorganic ana-lytical scheme. Such a precipitation will not, however, result in a film on a sub-strate or on the walls of the reaction vessel (actually, it may do so to a very slightextent but this film would be extremely thin). To form a visible film of CdS, con-ditions must be chosen so that bulk precipitation is prevented or at least sloweddown drastically. This is the purpose of the chalcogenide precursors, discussed inthe previous section. They slowly generate the chalcogenide, allowing slow for-mation of the metal chalcogenide (CdS in the present example).

The formation of the film, based on the formation of chalcogenide ions, canoccur by two fundamentally different processes. We continue to use CdS as theexample.

2.5.2 Ion-by-Ion Mechanism

The simplest mechanism, often assumed to be the operative one in general, iscommonly called the ion-by-ion mechanism, since it occurs by sequential ionic re-actions. The basis of this mechanism, illustrated for CdS, is given by

Cd2� � S2� → CdS (2.10)

If the ion product [Cd2�][S2�] exceeds the solubility product, Ksp, of CdS(10�28; Table 1.1), then, neglecting kinetic problems of nucleation, CdS willform as a solid phase (see Chap. 1). If the reaction is carried out in alkaline so-lution (by far the most common case), then a complex is needed to keep themetal ion in solution and to prevent the hydroxide from precipitating out (butsee later). Since the decomposition of the chalcogenide precursor can be con-trolled over a very wide range (by temperature, pH, concentration), the rate ofCdS formation can likewise be well controlled. Of course, the CdS should forma film on the substrate and (at least ideally) not precipitate in the solution. This


aspect of the ion-by-ion mechanisms (and all other mechanisms) is treated inSection 2.6.

2.5.3 Hydroxide Cluster Mechanism

It was stated earlier that complexation of the Cd was necessary to preventCd(OH)2 precipitation. However, very often (more often than realized), Cd(OH)2

(or metal hydroxides in general) are important reaction intermediates in the CDprocess. If the complex concentration is not high enough to prevent completely theformation of Cd(OH)2, then a relatively small amount of Cd(OH)2 may be formed,not as a visible precipitate, but as a colloid. Since Cd(OH)2 is colorless and col-loids typically do not scatter light, unless they aggregate to a large extent (in whichcase a suspension is the result), this means that the Cd(OH)2 colloid may not bevisible to the eye.

The CdS is then formed by reaction of slowly generated S2� ion with theCd(OH)2:

Cd2� � 2OH� → Cd(OH)2 (2.11)

followed by

Cd(OH)2 � S2� → CdS � 2OH� (2.12)

Reaction (2.12) occurs because Ksp for CdS (10�28; Table 1.1) is much smallerthan that for Cd(OH)2 (2 � 10�14). Another way of looking at this is that the freeenergy of formation of CdS is more negative than that of Cd(OH)2. It has also beensuggested that the hydroxide cluster can act as a catalyst for thiourea decomposi-tion. In this case, sulphide formation will occur preferentially at the surface of thehydroxide rather than nucleate separately in the solution. Such a course is logicalbased on the previous discussion of the effect of metal cations on the equilibriumof thiourea decomposition.

It should be borne in mind that the mechanism may change in the course ofthe deposition. As the metal is depleted from solution, the complex:metal ratiowill increase and may pass the point where no solid hydroxide phase is present inthe solution. In this case, the ion-by-ion process will occur (initially in parallelwith the hydroxide mechanism, later maybe exclusively) if the conditions are suit-able.

2.5.4 Complex-Decomposition Mechanism

The chalcogenide precursors possess many talents. Apart from forming thechalcogenide ions, they also form complexes with metal ions. As noted at the be-ginning of this section, and ignoring the distinction between ion-by-ion and hy-droxide cluster mechanisms treated previously, CD processes can be divided ac-cording to two basic mechanisms: participation of free sulphide ions (the


commonly accepted mechanism in most cases, although this does not necessarilymean that it is always the correct one), and a pathway involving decomposition ofa metal/chalcogen-containing complex without formation of free sulphide. Un-fortunately, it is usually difficult to distinguish between these two processes, andthe former is assumed more out of inertia than because of any clear proof. In thespecific case of the CdS deposition using thiourea, a complex-decompositionmechanism has been proposed in a number of different investigations, based onkinetic studies of the film formation process. Here we can revert to intuition andsuggest that, in the case of strong complexation between the chalcogen compoundand the metal ion (e.g., as occurs between thiosulphate and Hg, Ag, and Cu), itmay seem more logical for the fairly weak SMS bond to break than the very strongmetal–chalcogen bond. To be fair, these very strongly complexed cations are alsothose whose chalcogenides have a very low solubility product, and therefore verylittle free sulphide would be needed to form those metal chalcogenides.

As for the free chalcogenide processes, the complex-decomposition mech-anism can occur either by an ion-by-ion (or molecule-by-molecule, since free ionsneed not be involved directly) pathway, e.g.,

[SO3MSMHgMSMSO3]2� � H2O → HgS � SO42� � 2H� � S2O3

2� (2.13)

(the molecule of HgS can interact with other HgS molecules to form clusters andeventually crystals. Of course, they may also redissolve. These aspects are treatedin the following section), or by a solid-phase intermediate, e.g.,

[Cd(OH)2]n � SC(NH2)2 → [Cd(OH)2]nMSMC(NH2)2 (2.14)

[Cd(OH)2]nMSMC(NH2)2 → [Cd(OH)2]n�1CdS � CN2H2 � 2H2O (2.15)

with eventual exchange of all the hydroxide in the Cd(OH)2 to CdS.

2.6 NUCLEATION, ADHESION, AND FILM GROWTH

Probably the least-known aspect of the CD process is what determines the nucle-ation on the substrate and the subsequent film growth. In considering this aspect,we will treat the ion-by-ion and hydroxide cluster mechanisms separately, al-though there will be many features in common. The principles discussed shouldbe the same for both the free chalcogenide and the complex-decomposition mech-anisms.

2.6.1 Ion-by-Ion Growth

For nucleation to occur homogeneously in a particle-free solution by the ion-by-ion process, supersaturation, usually a high degree of supersaturation, is typicallyrequired (Chap. 1, section 1.2.1). The presence of a surface (the substrate or the


walls of the reaction vessel) introduces a degree of heterogeneity that facilitatesnucleation. For this reason, depositions that proceed via the ion-by-ion processtend to occur mainly on the substrate or other surfaces, rather than involving alarge amount of precipitate typical of the hydroxide mechanism. The surface canbe considered a catalyst for the nucleation.

As discussed in Chapter 1, the most important force involved in adhesion ofthe deposit to a substrate in general is the van der Waals force of attraction. In theinitial stages of growth, there may be specific chemical interactions between thedeposit and substrate. For example, if gold is used as a substrate, S, Se, and manyof their compounds interact chemically with the gold to form S(Se)–Au bonds.This would promote good adhesion of the deposit to the gold. There could also bechemical and electrostatic interactions between surfaces of the individual crystals.For example, the positive S(Se) face of polar crystals could bind to the negativemetal face of an adjacent crystal if the relative orientations are suitable (in prac-tice, this will probably not occur, since the crystal faces will adsorb solutionspecies as they grow). However, the van der Waals interaction between the crys-tals in the strongly ionic solution is enough in most cases to ensure adhesion of thecrystals to one another.

The fact that reasonably adherent films can be grown on apparently unreac-tive substrates, such as plastics, and even on such an inert and hydrophobic mate-rial as Teflon suggests that while such specific interactions between the semicon-ductor and substrate may improve adhesion to the substrate, they are not essentialfor film formation.

Once nucleation has begun on a substrate (this usually includes the insidewalls of the reaction vessel), it generally becomes easier for the film to grow, sincedeposition usually occurs more readily on the nucleated surface than on the cleansurface. The crystals will continue to grow until blocked by some process, such assteric hindrance by nearby crystals or adsorption of surface-active substancesfrom the solution. The former is probably the dominant reason for growth termi-nation in most cases.

2.6.2 Hydroxide Mechanism

Nucleation of the chalcogenide is much simpler in this process, since a solidphase—the metal hydroxide (or other solid phase)—is already present and theprocess proceeds by a substitution reaction on that solid phase. In this case, theinitial step in the deposition is adhesion of the hydroxide to the substrate. Thishydroxide is then converted into, e.g., CdS, forming a primary deposit of CdSclusters. More Cd(OH)2 and, as the reaction proceeds, CdS and partially con-verted hydroxide diffuses/convects to the substrate, where it may stick, either touncovered substrate (in the early stages of deposition) or to already depositedmaterial. This is essentially the same process as aggregation, described in Chap-


ter 1 in the section on forces, and again is a consequence of van der Waals at-traction.

Since the initial nucleation of hydroxide occurs homogeneously in the solu-tion, the CdS also is formed homogeneously and therefore usually precipitates outin the solution to a large extent. This precipitation occurs if the isolated crystalsaggregate to a sufficient extent to form large flocs. Film formation occurs whenhigh-surface-energy particles (single nanocrystals or small aggregates) reach thesubstrate (or any other surface) before they precipitate out in the form of large ag-gregates. This aggregation and homogeneous precipitation can be minimized, insome cases even prevented, by judicious choice of deposition parameters. Thus,while extensive precipitation suggests a hydroxide mechanism (some precipita-tion can occur in the ion-by-ion process), its absence does not always mean thatthe ion-by-ion process occurs.

An expected difference between ion-by-ion and hydroxide (or any othercluster) mechanisms is that in the latter, since colloids from the solution stick tothe substrate surface, the crystal size is not expected to change greatly with filmthickness (it may increase to some extent, since the colloids themselves can growvia an ion-by-ion process on the crystals). For ion-by-ion growth, it is likely thatcrystal growth occurs on nucleii already present on the substrate, and thereforecrystal size can increase with increasing deposition.

The foregoing description assumes adsorption of colloidal metal hydroxidefrom solution onto the substrate as the primary nucleation step. However, hy-droxides can also adsorb on solid surfaces at pH values below that of bulk hy-

FIG. 2.1 Schematic diagram showing the probable steps involved in the ion-by-ionmechanism. A: Diffusion of Cd and S ions to the substrate. B: Nucleation of the Cd and Sions facilitated by the substrate to form CdS nucleii. C: Growth of the CdS nucleii by ad-sorption of Cd and S ions from solution and nucleation of new CdS crystals. D: Continuedgrowth of CdS crystals, which adhere to each other through van der Waals forces (possiblyalso chemical interactions).


FIG. 2.2 Schematic diagram showing the probable steps involved in the hydroxidemechanism. A: Diffusion of hydroxide colloidal particles to the substrate, where they ad-here (B) and react with S ions (either generated homogeneously in solution or catalyzed bythe hydroxide surface). This reaction results in exchange of the hydroxide by sulphide,probably starting at the surface of the colloid and proceeding inward (C). This reaction willoccur both at the surface-adsorbed colloids and at those dispersed in the solution. Reactionwill continue (as long as the supply of sulphide continues) until most of the hydroxide isconverted to sulphide (D); eventually the primary particles of CdS will adhere to each otherto form an aggregated film (E); usually the nonadsorbed particles will also aggregate andprecipitate out of the solution.

droxide precipitation, an effect which has been related to the presence of an elec-tric field at the substrate/solution interface. This will certainly affect the nucle-ation, since this can now occur only on the substrate and not in solution. It is notso clear whether it will affect further crystal growth or not.

The basic features of the ion-by-ion and hydroxide cluster film-formingmechanisms are shown schematically in Figures 2.1 and 2.2, respectively. Filmformation involving complex decomposition will proceed in a similar manner(Fig. 2.3 shows this for a molecule-by-molecule deposition).


2.7 KINETICS OF THE DEPOSITION

Due to the different pathways that can occur in the CD process, the kinetics canvary widely from one deposition to another, as reflected by the rather wide rangeof activation energies (not commonly measured, but measured often enough todraw some conclusions) found. Regarding the time taken for a deposition, somedepositions can be completed in a few minutes or less, while others can proceedfor days and still be far from termination. This section is only meant to give a gen-eral picture; Chapter 3 should be consulted for more specific details and examples.

Kinetic studies on the growth of CD films show, in most cases, an inductionperiod at the beginning of the process where no clearly observable growth occurs,an approximately linear growth region, and a termination step where no furthergrowth occurs (see Fig. 2.4). Strangely enough, this type of growth kinetics oftenoccurs regardless of the deposition mechanism. For the ion-by-ion growth, it isvery simple to explain. Deposition begins only when the chalcogenide concentra-tion is high enough to allow nucleation to occur—the induction time correspondsto this buildup of chalcogenide concentration. Growth then occurs on these initialnucleii, along with new nucleation—the approximately linear region of growth.As the limiting reactant is used up, growth will start to slow down and eventuallystop due to depletion of the reactants.

For the cluster mechanism, while growth and termination can be similarlyexplained, the induction period is less obvious. The hydroxide cluster can start toadsorb on the substrate immediately after immersion of the substrate in the depo-sition solution, yet experiments have shown that film growth often does not occurfor some time. While the reason for this is not clear, it may be connected with the

FIG. 2.3 Schematic diagram illustrating possible steps in the complex-decompositionmechanism. The complex (CdMSML, where L is a ligand or part of the S-forming species)decomposes to CdS on the substrate (possibly catalyzed by the substrate) and, to a greateror lesser extent, also homogeneously in the solution (A, B). The CdS nuclei formed growby adsorption and decomposition of more complex species (C) until a film of aggregatedcrystals is formed (D) in the same manner as for the previous two mechanisms.


fact that the hydroxide particles often do not form a film, beyond some primaryadsorption on a surface; only when reaction to form the metal chalcogenide occursdoes film formation develop. In this case, we can again invoke the need for a min-imum concentration of chalcogenide ion. Some studies of the deposition rate havesuggested that the rate-limiting step is a chemical rather than a diffusion process,which supports the formation of the chalcogenide as this limiting step rather thandiffusion of cluster species to the substrate. Also, as described in the previous sec-tion, metal hydroxide might deposit on the substrate under conditions where it willnot form in the solution, but this is likely to be confined to the surface layer. Itmust be stressed, however, that due to the various possible processes involved inCD, the results of one or even of several studies cannot automatically be extrapo-lated to all other depositions of the same compound.

If reaction is allowed to proceed until the termination stage is reached, theterminal thickness of many CD films is typically several hundred nanometers, al-though it may reach a micron or more in some cases. This terminal thickness de-pends to a large extent on the deposition parameters. To take an extreme case, ad-dition of sulphide to a solution of Cd ions will give an immediate precipitate ofCdS, but no (or at most an extremely thin) deposit on the walls of the depositionvessel, which may thicken somewhat with time, but will not be visible (whichmeans a terminal thickness less than ca. 20 nm). For a normal CD reaction, if pre-cipitation occurs homogeneously in solution, then that precipitate is lost for film

FIG. 2.4 Typical shape of the curve reprenting time dependence of film thickness dur-ing growth.


deposition, with resulting reduction in terminal thickness. Therefore, the ion-by-ion process, with its lesser tendency for homogeneous precipitation, will usuallyresult in a larger terminal thickness than the cluster process, for comparable initialreactant concentrations. Of course, film thickness can made as small as desiredsimply by removing the substrate when the desired thickness has been reached (al-though very thin films may not be homogeneous but, rather, clusters spread het-erogeneously on the substrate). Alternatively, films thicker than the terminalthickness may be obtained by repeated deposition (there will be a limit to thethickness even here, since thick films tend to peel off the substrate).

2.8 SUBSTRATE

To a first approximation, films can be deposited by CD on any surface (this is oneof the advantages of CD). Of course, there will be certain obvious exceptions, suchas substrates that are unstable in the deposition solution (this is rarely a problem,in practice) or “dirty” substrates. In several studies, CdS has been shown to formquite adherent films on Teflon, and this attests to the ability of CD to form films ona wide range of substrates. An important advantage of CD is that the shape of thesubstrate is usually not important – very irregularly-shaped substrates can be used.

To a second approximation, the nature of the substrate is usually importantin order to obtain an adherent film; some substrates result in more adherent filmsthan others. Rough substrates are better in this respect (in common with most de-position procedures), probably due to the greater actual surface area of contact pergeometric surface area and the possibility of anchoring of the initial deposit inpores of the substrate. Oxides [this includes glass, conducting oxides such as tinoxide and indium tin oxide, and, to a lesser extent, silica (quartz)], in spite of ap-pearing inert, are actually quite reactive in terms of their adsorption properties.This is due to the presence of hydroxyl surface groups, which can form fairlystrong hydrogen bonds. Yet there can be very noticeable differences in the adhe-sion to different glasses and between glass and silica (deposits tend to be less ad-herent to silica than to glasses in general). An early study on CD PbS films usingdifferent glass substrates found large differences in film formation on the differ-ent substrates; no (or, at best, only patchy) films were formed on borosilicate glassor on silica, whereas lead flint glasses, followed by zinc crown glass, resulted inthe best films [25]. The ability to form good PbS films on these latter glasses wasascribed to the ability of the PbO or ZnO in these glasses to from insoluble sul-phides. This would enhance binding to the depositing film. The possibility of ionexchange between metal ions in the glass and those in solution may also play arole in binding the initial CD film.

Glass substrates can be sensitized, usually with a solution of SnCl2, whichhydrolyzes to give nuclei of tin hydroxide or oxide, on the surface. While in mostcases of CD, such sensitization is not used, and not required, there have been re-


ports of better layers (more adherent and/or homogeneous layers, faster deposi-tion) using such sensitization. For example, it has been shown that, for PbSe de-position from selenosulphate solutions, film formation begins immediately onSnCl2-sensitized glass, instead of after an induction period as on plain glass. Ad-ditionally, film formation occurs (at least initially) in the absence of bulk precipi-tation in solution, in contrast to the parallel film deposition and bulk precipitationthat occurs when nonsensitized glass is used [26]. This suggests that a high su-persaturation is required for nucleation to occur on untreated glass (as well in so-lution) and that this is not the case when nuclei are already present on the glassfrom the sensitization (or from previously deposited PbSe). It can be concludedthat such sensitization of glass (and probably many other surfaces) should be con-sidered if satisfactory growth is not obtained without it. Although probably notimportant for most requirements, it should also be kept in mind that such films willcontain a small amount of tin at the film/substrate interface.

Metals make good substrates in general, either because chalcogenides tendto adsorb strongly on the noble metals, in particular gold, or the non-noble metalsare covered with a (usually hydroxylated in the deposition solution) oxide layer.In addition, if the metal in the deposition solution has a sufficiently negative po-tential, an internal electrochemical reduction may occur (remember that elec-trodeposition can often be carried out from CD solutions). This was suggested along time ago for deposition of various metal sulphides from thiosulphate solu-tions on certain metal substrates [3,5].

A large variety of CD films have been deposited on different polymer sur-faces subjected to various activation treatments [27]. The most effective treatmentwas immersing the substrate in KMnO4 solution for 24 hr, which formed a brownMn-O film and subsequent removal of this film with, e.g., conc. HCl. It was sug-gested that the permanganate introduces carboxylic groups on the originally hy-drophobic surface.

Films have also been deposited on monolayers—both Langmuir and self-as-sembled. In many cases, such depositions have studied nucleation of very thin lay-ers of deposit (more accurately, scattered nanocrystals), but deposition of thickerfilms, such as PbS [28] and Fe(O)OH [29] on self-assembled monolayers and var-ious metal sulphides and selenides on Langmuir–Blodgett films [30] have beenstudied, in some cases to understand how a well-defined—either chemicallyand/or geometrically—substrate can control nucleation and growth geometry. De-position does not occur on some monolayers, or at least considerably less readilythan on non-monolayer-covered substrate. An example of such a monolayer is oc-tadecylphosphonic acid; CdS was found not to deposit on this monolayer but didgrow on the free areas of mica partially covered with the monolayer [31].

The use of monolayers as substrates has been exploited to pattern deposits.The principle behind this idea is that monolayers with either hydrophobic orhydrophilic end groups can be patterned onto a substrate. Deposition will usually


occur only (or at least highly preferentially) on the hydrophilic endgroups.Whether this is due to a simple physical interaction between solution (aqueous)species and a hydrophilic surface, due to electrical charge on the solvated end-group, or to some other specific interaction (or some combination of these effects)is not clear. An example of this patterning is shown here for TiO2 deposition [32].A long-chain thioacetate-terminated trichlorosilane is self-assembled on anoxidized Si substrate (Fig. 2.5a and b). The thioacetate is somewhat hydrophobic.Exposure of this monolayer through a mask (a fine mesh grid) to UV radiation

FIG. 2.5 Self-assembly of a long-chain thioacetate-terminated trichlorosilane on an ox-idized Si substrate (a and b). Exposure of this monolayer through a fine mesh grid mask toUV radiation (c) oxidizes the somewhat hydrophobic-thioacetate to hydrophilic sulphonateendgroups (d). Deposition of TiO2 from an aqueous solution of TiCl4/HCl at 80°C on theUV-exposed (hydrophilic) regions of the substrate (e). (From Ref. 32.)


oxidizes the thioacetate to hydrophilic sulphonate groups (c and d). Deposition ofTiO2 from an aqueous solution of TiCl4/HCl at 80°C occurred only on the UV-ex-posed (hydrophilic) regions of the substrate (Fig. 2.6). See Section 4.1.5.3 andSection 5.2.4.2) for other examples of patterned deposition (of CdS and PbS,respectively) using different patterning procedures.

For any particular substrate, the adhesion can also depend on the depositionparameters, although little is reported on this aspect (for a good reason—it is notreally understood). The surface of glass (of whatever type) is hydroxylated inaqueous solution and the concentration of the various species is clearly pH de-pendent, as can be seen from Eq. (2.16):

(2.16)

[Since silica is very acidic, the dissociation to give a positively charged surface(MSiMOH2

�) will only occur in very acidic solution (pH � 2) and is not com-monly encountered in CD. Other less acidic oxide surfaces may, however, be pos-itively charged in CD solutions.] This dissociation can affect interaction betweenglass and various species in the solution as a function of the solution pH. Note thatwater (also ammonia) may interact, through hydrogen bonding, with the hydrox-ylated silica.

In the end, adhesion as a function of deposition parameters (in contrast tothe nature of the substrate) can usually not be reliably predicted, and there is nosubstitute for experimental experience.

Apart from adhesion, the crystallographic properties of the CD film aresometimes dependent on the nature of the substrate (although more often theredoes not seem to be any dependence of this type). One example is epitaxial depo-sition on a crystallographically ordered substrate [epitaxial here means a struc-

SiMOMHMJJ SiMO� � H�M

JJD

FIG. 2.6 SEM micrograph of patterned TiO2 deposit (left side) and Ti elemental map-ping of this sample (right side). (From Ref. 32.)


tural relationship between the crystal lattice of the deposit and of the substrate—i.e., the crystal axes of the deposit are aligned (but not necessarily parallel to) thecrystal axes of the substrate]. Examples of this will be discussed later and in otherchapters. Also, even if the deposition is not epitaxial (and only a few cases havebeen reported where epitaxy occurs), different texturing of the film may occur.Texturing refers to the preferred orientation of the crystals of the deposit perpen-dicular to the substrate; thus, textured films of hexagonal CdS often have the basalface (the Cd or S face) pointing upward. If orientation occurs also in the plane ofthe substrate (azimuthal orientation), then the deposit is called oriented. The termorientation is often (erroneously) used where texturing is meant.

A monolayer-covered substrate may dictate the crystallographic form of aCD film, depending on the interaction between the monolayer endgroups and thesemiconductor. Thus, epitaxial growth of PbS has been accomplished at arachidicacid monolayers where the PbS interatomic spacings along the (111) plane arewell matched with the monolayer packing [33]. Texturing has been observed insome cases, even when the monolayers themselves were not well ordered.

It will be obvious that the cluster mechanism of deposition is unlikely tolead to an oriented film, since the clusters would have to align themselves with thesubstrate lattice, either on adsorption or subsequently. Therefore an epitaxial filmis highly suggestive of an ion-by-ion growth, which is more likely to be directedby the substrate.

One example has been described of CD (PbS) on a poled ferroelectric sub-strate. The PbS crystal size was larger (ca. 1 �m) on the poled substrate than onthe unpoled (or a glass) substrate (ca. 0.3 �m) [34]. Other changes in the electri-cal properties of the films were noted. The differences were ascribed to the elec-tric field and charge accumulation at the ferroelectric surface (more details can befound in Sec. 5.2.4.3).

2.9 DEPOSITION OF SPECIFIC SEMICONDUCTORS

This section will give an overview of the various semiconductors that have beendeposited by CD. The groups of semiconductors will be divided in much the sameway as in the rest of the book:

II–VIIV–VIOther binary sulphides and selenidesOxides, other binary semiconductors, and elemental SeTernary compounds

A few specific examples of experimental details will be given in this sec-tion, with explanation of the importance of the different variables.


2.9.1 II–VI Compounds

This comprises the most studied group and includes CdS (the single most studiedcompound), CdSe, CdTe, ZnS, ZnSe, HgS, HgSe, and various mixtures of thesecompounds.

2.9.1.1 CdS

Film Deposition. We begin with a common “recipe” for CdS depositionand follow the deposition process using this recipe. This is described in detail asa way of understanding both the deposition process and which factors can affectthe final film.

STEP 1: PREPARATION OF STOCK SOLUTIONS. Usually, stock solutions of thesolid reactants are prepared in advance, since many depositions are normally car-ried out. The three reagents required are:

CdSO4: Other water-soluble Cd salts can be used equally well, such as thechloride or acetate; there is no clear evidence in general that the nature ofthe anion is important, although there are a few studies that reported dif-ferences in the deposit depending on the anion used.

NH4OH: Since this is required in high concentrations and is a liquid, a stocksolution is not required, but it can be used directly from a bottle of con-centrated ammonia (concentrated ammonia is ca. 15 M).

Thiourea (SC(NH2)2): This solution will slowly precipitate sulphur (seenas a fine white precipitate—white instead of the usual yellow of sulphurpossibly due to size quantization of the finely divided precipitate?).However, it can usually be kept over a period of weeks or even monthsin a stoppered bottle without major adverse effects, although for opti-mum reproducibility a fresh solution may be preferred. If a preparedstock solution is used, filtration of this solution before use will mini-mize the presence of sulphur particles in the deposition solution; suchparticles can act as nucleation centers and accelerate precipitation in thesolution.

STEP 2: PREPARATION OF THE SUBSTRATE(S). While almost any substrate canbe used, we will use glass microscope slides in this example; this is a commonsubstrate and makes it easy to see the CdS film. The microscope slide can be cutto whatever size and shape is convenient. The slide should be cleaned well, sincefilms usually do not adhere well to “dirty” surfaces. Suitable cleaning agents aretrichloroethylene or/and sulphochromic acid, and the slide should be well rinsedwith pure water. If the slide is clean, water dropped onto it will form a film (hy-drophilic surface), while on a “dirty” (hydrophobic) slide the water will formdrops. Needless to say, the part of the slide where deposition is to occur should notbe touched with the hands after this treatment.


STEP 3: PREPARATION OF THE DEPOSITION SOLUTION. Solution compositions,even based on exactly the same constituents as stated earlier, vary widely from oneliterature source to another. We give a typical “average” composition and discussthe effect of variations from this composition.

This “average” solution is made up of (concentrations in final solution):

CdSO4 10 mMAqueous ammonia 1MThiourea 50 mMSolution pH 11Deposition temperature 70°C

Concentrated ammonium hydroxide is added to a stock solution of CdSO4 (orother Cd salt). Initially, Cd(OH)2 precipitates, but this redissolves in excess am-monia to give the cadmium ammine complex:

Cd2� � 4NH3D Cd(NH3)42� (2.17)

This solution is heated to 70°C in a thermostated water bath. A stock solution ofthiourea is then added to bring the thiourea concentration up to 50 mM. The finalsolution pH can be adjusted with KOH (more basic) or acetic acid (more acidic).

STEP 4: DEPOSITION. The substrate(s) is immersed in the preceding solutionand placed in the water bath at 70°C. (Note that CdS usually also forms homoge-neously in solution, and this can sediment onto the substrate, where it forms aloose coating and may prevent growth of an adherent film. For this reason, thesubstrate should be placed vertically in solution; or if placed at an angle or evenhorizontally, only the underside of the substrate, where sedimentation does not oc-cur, should be used.)

The time of deposition is variable (although it should be reproducible fromone run to another under the same conditions), and it is difficult to give a particu-lar time. Usually a thin film (tens of nanometers) will form in some minutes, andthis will slowly thicken over some tens of minutes to hours to a typical terminalthickness of ca. 200 nm. The simplest way of determining the optimum time (and“optimum” will depend on the application) is to simply look at the film. With a lit-tle experience (or with the more quantitative help of a spectrophotometer), the ap-proximate thickness can be estimated from how deep the yellow color is. Thick-ness can also be estimated from the transmission spectrum by measuring the dropin transmission over the near-bandgap region; this method is less dependent on re-flection and scattering losses (as long as they are not too large) than simply mea-suring absorption at a fixed wavelength. In any case, prior calibration of the spec-tra with known thicknesses of the film is required.

When the desired thickness is attained, the substrate is removed and rinsedwith water. This rinsing can also be carried out in an ultrasonic bath, which moreeffectively removes loose particles.


VARIATIONS IN THE PROCESS. There are many variables in this (experimen-tally simple) process; concentration of the various reactants, pH, and temperatureare the main ones. Other, less important (in most cases) ones are stirring of solu-tion and illumination of the solution during deposition. We will ignore these lasttwo here.

The effect of reactant concentration can be divided into two separate influ-ences. The simplest is obvious: Lower overall concentrations result in a slower rate.This does not necessarily mean a thinner film, however—sometimes the opposite.The reason for this is clear if we return to our introductory discussion on the CDprocess—rapid precipitation. It is clear that if the reaction is too fast, it will termi-nate with most of the product precipitating homogeneously in solution rather thandepositing on the substrate (which requires time to occur). This results in a verythin film, if any film at all. Similarly, for the less extreme case of a CD reaction thatterminates, not within a second, but still in a short time, the final film thickness willbe small. At the other extreme, if the reaction is extremely slow, a thick film can bebuilt up, but it may take a very long time for this to occur (weeks, even months). Itis therefore evident that there is an optimum rate for the reaction, which can be con-trolled by a combination of reactant concentrations, temperature, and pH.

A separate effect of concentration is the ratio between the metal ion and thecomplexant concentrations. This ratio determines, often more than the overall re-actant concentrations, the reaction rate, since it controls the concentration of freemetal ions in solution. It can also determine the reaction pathway. Further discus-sion of this factor will be left for the next example (CdSe), since it has been treatedin more detail for that case.

The solution pH influences a number of factors, and it is not always simpleto predict its effect. Thus, thiourea decomposition (in alkaline solution) is gener-ally faster at higher pH. The probability of the presence of a solid phase ofCd(OH)2 and its concentration in the solution are both increased at higher pH(higher OH� concentration). The pH is determined, in the example given earlier,by the concentration of ammonia. However, it can be adjusted independent of theammonia concentration. Addition of an ammonium salt, which with ammonia,acts as a buffer, will lower the pH through the following equilibrium:

NH3 � H2OD NH4� � OH� (2.18)

Ammonia is alkaline in water because of this equilibrium, which produces hy-droxide ions. The addition of an ammonium salt (NH4

� ions) will push the equi-librium back toward the left, i.e., lower hydroxide concentration, therefore lowerpH. Increase of pH can be effected by addition of sodium or potassium hydroxide.This explains why some CdS depositions based on the Cd/NH3/thiourea formulainclude either an ammonium salt or an alkali metal hydroxide.

Finally, the effect of temperature on increasing the reaction rate (possiblyalso the mechanism—see next example) is again obvious, since the thiourea de-


composition will be faster at higher temperature. Additionally, since the stabilityconstant of the complex is usually smaller at higher temperatures (see Chap. 1),there will be a higher free metal ion concentration, which again translates intofaster rate (although, since sulphide generation is usually rate determining, this ef-fect may not be large). Temperature programming during the deposition has alsobeen employed. This can be useful, for example, if the deposition solution changesthe substrate in some beneficial way (as occurs in some photovoltaic cell sub-strates—see Chap. 9). The initial low temperature delays coverage of the substrateand allows more time for this surface treatment to occur. If the temperature is thenincreased, deposition occurs at a reasonable rate.

There are other bath compositions based on different sulphide-generatingprecursors and/or complexing agents. Thioacetamide and thiosulphate are two ofthe former, while ethylenediamine is a common example of a complexant that hasbeen used instead of ammonia. The volatility of ammonia, and its gradual loss inan open deposition bath, is circumvented by using a less volatile complexant, suchas ethylenediamine.

Some Properties of CdS Films. It should be noted that all properties dis-cussed in this chapter refer to as-deposited films (not annealed) unless specificallystated otherwise. In general, annealing increases crystal size and reduces dark con-ductivity. The latter obviously depends to a large degree on the annealing atmo-sphere.

STRUCTURE. CdS can exist in three different crystal structures: hexagonal(wurtzite), cubic (zincblende)—both tetrahedrally coordinated and cubic (rock-salt), which is sixfold coordinated. Except in a few cases, the rocksalt modifica-tion of CdS has been observed only at very high pressures: CD films of this phasehave never been reported. The other two phases have been reported to occur in CDfilms under various conditions. The wurtzite phase is thermodynamically slightlymore stable, and invariably forms if the zincblende phase is heated above300–400°C. The low-temperature CD method therefore can allow the formationof the zincblende phase, and this phase is commonly obtained in CD CdS films.Very often, a mixture of wurtzite and zincblende phases has been reported in theliterature. There are many variables that affect the crystal structure, including thenature of the complex, the substrate, and sometimes even stirring.

ORIENTATION. In common with with other deposition techniques, a prefer-ential texturing of the film in the (111) (for zincblende) or (0002) (for wurtzite) di-rection is often reported. However, nontextured (or weakly textured) films areprobably more commonly obtained. As for crystal structure, the degree of textur-ing depends on several factors, an important one being the nature of the substrate.It is likely that highly textured deposits form by the ion-by-ion mechanism; a puresolid-phase cluster mechanism is less likely to result in strong texturing.


In certain cases, epitaxial growth has been observed on crystallographicallywell-defined substrates, such as single crystal InP and CuInSe2. This type ofgrowth, even more than textured growth, demands ion-by-ion growth as well as acrystallographically well-defined substrate. Lattice parameters (at least at somedefined ratio and angular match) fairly close to that of CdS are also required for arelatively defect-free interface, although there is some flexibility here. For exam-ple, CdS with a high density of stacking faults has been epitaxially deposited onGaP (7% mismatch).

OPTICAL PROPERTIES. The most commonly reported optical properties areoptical transmission, with some studies also on photoluminescence. The impor-tance of the optical transmission for CdS in particular lies in its use in photovoltaiccells, where it acts as a window layer. The CdS should be as transparent as possi-ble to the incoming radiation. The transmission is a function of thickness,bandgap, and film structure (is the film transparent or scattering?). The bandgapin most studies is constant (ca. 2.45 eV at room temperature), although somewhatlarger values have been obtained due to size quantization in very small crystallites.

Photoluminescence of the films varies greatly, both in intensity and in spec-tral shape, from one report to another. This is not surprising, since this property isvery dependent on the state of the surface of the individual crystals. A red (ca. 1.8eV) defect emission is usually seen, but green, yellow, and infrared peaks havealso been reported. The various wavelengths are related to different defects in/onthe crystals; even the green emission is probably due to a shallow defect emission.

ELECTRICAL PROPERTIES AND PHOTOCONDUCTIVITY. Electrical resistivity ofCD CdS films is commonly studied. Values for this (dark) resistivity vary overmany orders of magnitude from one film to another, usually for reasons that are notunderstood. Values as high as 109 �-cm and as low as 15 �-cm have been reportedfor undoped films (doped films have been reported with still lower resistivities).

Since the films are often highly resistive, it is not surprising that they exhibitstrong photoconduction. Photoconduction occurs due to the formation of free car-riers by illumination, and if the free carrier concentration is low to begin with (lowconductivity), then the photogenerated carriers will usually dominate the conduc-tivity. This is in contrast to a relatively conducting semiconductor (high dopinglevel), when the extra photogenerated (majority) carriers are only a small pertur-bation to those present in the dark. Light:dark photoconductivity ratios (sensitiv-ity) as high as 109 have been reported. In many cases, the photocurrent decay timeis measured in hours, and this is explained (in a general way) by slow states; thenature of these states is not usually known.

2.9.1.2 CdSe

Film Deposition. Chemical deposition films of CdSe have been relativelywidely investigated, largely for photoanodes in photoelectrochemical cells (see


Sec. 2.12.1.2). Here we will give an example of CdSe deposition, emphasizing thedifferences between the procedures for CdSe and CdS. The most important singledifference in the deposition is, of course, the chalcogenide-forming reagent. Aprocedure used commonly by the author will be described, where the complexantis nitrilotriacetic acid (NTA), which is related to the more common EDTA(ethylenediaminetetracetic acid) and has the chemical formula N(CH2COOH)3. Itis a strong complexant for Cd2� (and many other cations), although less so thanEDTA, but stronger than the often-used ammonia. It also has an advantage (inmost cases) over ammonia in that it is nonvolatile and is not lost during deposi-tion.

STEP 1: PREPARATION OF STOCK SOLUTIONS. The aqueous Cd solution is pre-pared as for CdS. As for CdS, the nature of the anion does not appear to be veryimportant—the sulphate is commonly used.

NTA. NTA itself is not very soluble in water. Either it is used as the Na orK salt, or the acid can be dissolved together with three equivalents ofKOH or NaOH to form a solution of the salt. This salt will be referred tohereafter simply as NTA.

Sodium selenosulphate (Na2SeSO3). This compound is made by dissolvingelemental Se in an aqueous solution of sodium sulphite:

Na2SO3 � SeD Na2SeSO3 (2.19)

The prepared solution will slowly deposit Se as a black precipitate over a periodof weeks. Also, the solution is considerably less stable than thiourea. If kept outof excessive contact with air, it will be usable for about a month if high repro-ducibility of the deposition kinetics is not important. However, it is important tobe aware of the slow decomposition and loss of reactivity of this reactant. If afreshly made Na2SeSO3 solution is used to deposit CdSe, the reaction will proceedmuch faster than if an aged solution is used; this fact should be taken into accountin preparing the overall deposition solution.

The solution is made up from an aqueous solution 0.4 M in Na2SO3 and 0.2M in Se. This solution typically is stirred at 60°C for a couple of hours (the Se dis-solves slowly, the reason a fresh solution is usually not made up every time). Ifthis solution is prepared over a longer time (maybe 6 hours and/or at a higher tem-perature, even boiling), it will undergo accelerated aging. This results in more re-producible deposition conditions on one hand, but at the cost of reduced reactiv-ity and a shorter lifetime of the solution.

STEP 2: PREPARATION OF THE SUBSTRATE(S). The same as described earlierfor CdS.

STEP 3: PREPARATION OF THE DEPOSITION SOLUTION. As for CdS, the condi-tions can be very variable. The following is a typical deposition solution, but re-


member that allowances should be made for the varying activity of the selenosul-phate. If deposition occurs too rapidly (slowly), parameters should be changed toslow down (speed up) the reaction [e.g., lower (higher) selenosulphate concentra-tion, higher (lower) NTA:Cd ratio, lower (higher) temperature].

The following procedure is followed to obtain a final concentration of 60mM each Cd2� and Na2SeSO3 and ca. 100 mM NTA (the reason for the “ca.” willbe explained shortly).

Take the calculated amount of stock Cd solution and add water. Then add thestock NTA solution. If the concentrated Cd and NTA solutions are mixed togetherwithout adding water between them, they may form a gel and the solution will thenbe useless. The pH should then be adjusted to ca. 8 by adding KOH or NaOH. It isimportant that the pH at this stage be greater than 7, since selenosulphate immedi-ately decomposes to red selenium at a pH lower than 7. This solution is brought tothe temperature at which the deposition is to be carried out. The selenosulphate(also at the same temperature) is then added. This will increase the pH (since theselenosulphate is alkaline). KOH is added as necessary to bring the pH to ca. 10(9.5–10 is a suitable range; it should be neither much lower nor much higher thanthis). If the pH increases to much above 10 on the addition of selenosulphate, di-lute acetic acid can be carefully added to lower the pH to the required range.

STEP 4: DEPOSITION. This is the same as for CdS, except that the depositionusually is slower and takes from hours to days, depending on conditions. Or-ange/red coloration in the solution, corresponding to the start of CdSe formation,should occur some minutes after adding the selenosulphate. If it occurs immedi-ately or almost immediately, the reaction may be too fast and only a thin terminalthickness may be obtained. On the other hand, if coloration has still not started af-ter about 30 min, the deposition will probably be very slow.

VARIATIONS IN THE PROCESS. The reader may have noticed that some of thepreparation details just given are rather vague. This reflects the varying activity ofthe selenosulphate solution, both with time and from one batch to another. It ismore important to understand the effects of the various parameters and to be ableto vary them logically than to follow an exact recipe. These effects will be de-scribed in detail in Chapter 3, and it is recommended that anyone wanting to de-posit these films read that chapter before carrying out the deposition. Here, a briefexplanation of these various factors is given.

The deposition temperature obviously will increase the reaction rate (andthe deposition rate, although, again, if the reaction is too rapid, only a thin filmwill be obtained). Another effect of increased temperature is increased crystal sizein the film. The crystal size varies (assuming the NTA:Cd ratio is not too high),typically from �4 nm at close to 0°C to ca. 8 nm at 80°C. This change can be seenby a change in color from yellow-orange (low temperature) to red (high tempera-ture) due to size quantization (see Sec. 2.12.2 and Chapter 10).


If crystal size is not an important issue, then higher temperatures are usuallymore convenient, since the deposition is faster. Adhesion of the film to the sub-strate also appears to be better in general at higher temperatures (larger crystalsize? See table 1.3).

Solution composition is important not only because the reaction rate in-creases with the concentrations of selenosulphate and/or Cd, but even more sothrough the ratio between the NTA and Cd concentrations. The higher this ratio,the slower the reaction, since the free Cd2� concentration is less. The optimumvalue for this ratio will depend, among other factors, on the deposition tempera-ture. At higher temperatures, a higher NTA:Cd ratio is required to prevent toorapid reaction. Thus the 100 mM concentration of NTA (for 60 mM Cd) given ear-lier is typical for ca. room temperature deposition, but would probably need to beincreased to 110 mM or even more at substantially higher temperatures.

Even more important, if the NTA:Cd ratio increases above a certain value(which depends on the other solution parameters but which is typically between1.7 and 2.1), the mechanism of deposition changes from a hydroxide cluster mech-anism to ion-by-ion deposition. The latter is considerably slower, does not deposithomogeneously in solution (or much less than the cluster mechanism), and resultsin larger crystal size (typically 8–20 nm, depending on deposition temperature).Also, in contrast to the cluster mechanism, where deposition is (should be) homo-geneous and the films transparent throughout the deposition, the ion-by-ion mech-anism usually results in films that appear nonhomogeneous and highly scatteringin the early stages of deposition, but become homogeneous and transparent withtime (typically several days at lower temperatures, less at higher ones).

If all the foregoing gives the reader the impression that CdSe deposition ismore of an art than a science, this would not be a gross misunderstanding. Expe-rience is certainly useful here, more so than for CdS deposition, which is more re-producible, probably due to the more stable thiourea (and possibly also because ofthe shorter deposition time typical for the CdS deposition, based on one of Mur-phy’s laws—the more time you allow for something to occur, the greater the op-portunity that something will go wrong).

Some Properties of CdSe Films.

STRUCTURE. CdSe forms the same three crystal structures as described ear-lier for CdS. The main difference between the CD films of the two materials isthat, while CdS can be commonly found in both the wurtzite and sphalerite forms,CdSe is more commonly deposited in the cubic zincblende form. Mixtures of thetwo forms have been reported in some cases, particularly when a visible Cd(OH)2

precipitate is present in the initial deposition solution.As for CdS, CdSe has been epitaxially deposited on single crystal InP. As

expected, epitaxy occurred only for the ion-by-ion mechanism, where individualspecies could either adsorb on or migrate to the ideal lattice position.


The CdSe crystal size in the hydroxide mechanism is probably determinedby the size of the CdOH)2 particles in the solution and on the substrate, while thatformed through the ion-by-ion mechanism will depend on the heterogeneous nu-cleation on the substrate, and is invariably larger.

OPTICAL PROPERTIES. CD CdSe films often exhibit size quantization, witha blue shift in their absorption spectrum of as much as 0.6 eV.

Photoluminescence of the films (in the absence of water vapor) usually isdominated, as for CdS, by a broad defect emission varying from ca. 1.4 to 1.6 eV.A (close to) band-to-band emission is often also observed, usually (but not al-ways) at a lower intensity than the broad defect emission. In the presence of wa-ter vapor, however, the band-to-band emission often dominates.

ELECTRICAL PROPERTIES. Resistivity studies on CdSe are much lesswidespread than on CdS films. The dark conductivity of undoped films is high(108 �-cm is typical), and the photocurrent sensitivity is less than for CdS films(even under illumination, the films are normally very resistive).

2.9.1.3 CdTe

Only two different studies on true CdTe deposition by CD appear to exist. Inthe first, CdTe was deposited from solutions containing TeO2 and hydrazine,the latter presumably slowly reducing the TeO2 to telluride ion. In this study, themain interest was not on the deposition itself but on the further use of the CdTefilms, and not much information was given on the films themselves. Morerecently, CdTe was deposited using what was apparently the Te analogue ofselenosulphate, Na2TeSO3. In all these depositions, the films were not verystoichiometric and included considerable amounts of Te. Some structural, optical,and electrical properties of these latter films were given (described in Chap. 4).In a very recent study, CD Cd(OH)2 films were converted to CdTe using asolution of Te in hydroxymethane sulphinic acid, which acted as a telluridesource. The conversion of the hydroxide to CdTe was incomplete, but theredid not seem to be any free Te in the films. Properties of these films were alsodescribed.

2.9.1.4 ZnS and ZnSe

The most important difference between CD of CdS(Se) and ZnS(Se) is related tothe difference between the solubility products of the hydroxide and chalcogenidesof the two metals. Considering the sulphides, the various values of Ksp are:Cd(OH)2—2 � 10�14; CdS—10�28; Zn(OH)2—1.10�16; ZnS—3.10�25. The dif-ferences between the pairs of Ksp are 2 � 1014 (for Cd) and 3 � 108 (for Zn). Sincehydroxide ions are present at much higher concentrations than sulphide, hydrox-ide formation, and stability against sulphurization, is much more likely for Zn thanfor Cd.


ZnS is commonly deposited by CD. However, in many cases, the ZnS is notstoichiometric and contains hydroxide in one form or another. ZnSe, on the otherhand, is more difficult to deposit from selenosulphate baths, in spite of the some-what lower solubility product of ZnSe (4.10�26) compared with ZnS. In fact, astrong reducing agent such as hydrazine is normally required. The hydrazine pre-sumably reduces the selenosulphate to give a high enough selenide concentrationto allow ZnSe to form, although it is possible that other factors are also important.(The sulphite, present in excess in selenosulphate, is itself a reducing agent, al-though much weaker than hydrazine.) Activation energies for ZnS(Se) deposi-tions are generally considerably lower than for CdS and PbS deposition, and thissuggests a different deposition mechanism (although, as we have already seen, thedeposition mechanism is largely a function of the deposition conditions and notsimply the material itself). Also, the crystal size of ZnS(Se) is usually smaller thanfor CdS or PbS. Both these factors suggest that ZnS and ZnSe form by a pure clus-ter mechanism. In fact, nowhere in the literature is there evidence for ion-by-iongrowth of ZnS or ZnSe.

2.9.1.5 HgS and HgSe

There are only several reports on the deposition of the mercury chalcogenides(some ternaries containing Hg have also been described).

HgS films have been deposited from a simple chemical precipitation reactionbetween mercuric chloride and sodium sulphide. Under suitable conditions, a filmis formed along with the precipitate. The thickness of the films were ca. 0.7 �m,which is thicker than normal for CD films, even more so considering the rapid na-ture of the reaction, which normally only leads to a very thin film, if any at all. Thefilms apparently deposit from the colloidal HgS formed on mixing the reactants.

HgS films have also been deposited from “more conventional” depositionsolutions using thiourea and the tetraiodide complex of mercury—a strong com-plex—in alkaline solution. Both these and the previous films showed an opticalabsorption with a gradual absorption onset at 700 nm and a sharp one at 400 nm.

The thiosulphate mercury(II) complex in ammoniacal solution has also beenused to deposit films of HgS, both in ammoniacal and nonammoniacal solutions.The deposit from the former was predominantly �-HgS (cinnabar). As with the pre-vious films, a sharp optical absorption onset at ca. 400 nm was observed, togetherwith a more gradual one extending, in this case, to beyond 800 nm and dependenton film thickness. The nonammoniacal solution gave crystal sizes (and opticalbandgaps) that varied with deposition temperature from 3 nm (2.4 eV) at 0°C to 8nm (1.9 eV) at 85°C and corresponding resistivities between 104 and 103 �-cm.

HgSe was first deposited from an iodide-complexed solution using se-lenosemicarbazide. Two other depositions were described, both using selenosul-phate. One used an alkaline Hg–formamide complex. The as-deposited films didnot show an XRD pattern, suggesting that the material was either amorphous or


very small nanocrystalline. The latter in particular is supported by the largebandgap measured from the absorption spectrum (1.42 eV compared to semimetalbulk HgSe with a negative bandgap). Electrical conductivity measurements indi-cated a midgap Fermi-level characteristic of intrinsic semiconducting material. Inthe other selenosulphate method, HgSe films were deposited from an ammoniacalbath onto polyester substrates. The films were strongly (111) textured. The crys-tal size was 7.7 nm, resulting in a strong blue shift in the optical absorption spec-trum, and a measured bandgap of 2.5 eV (compared to a negative—semimetal—bandgap of bulk HgSe). The sheet resistance of these films (13 k�-cm�2) wasrelatively low, considering the small crystal size.

2.9.2 IV–VI Compounds

2.9.2.1 Deposition

As noted at the beginning of this chapter, most of the early studies in CD focusedon PbS, followed by PbSe, driven by their photoconducting properties. For opti-mum use as photoconductors, the deposited films were annealed in an oxygen-containing atmosphere. Most of this section will focus on nonannealed films, andannealed films will be treated only very briefly. More details on the annealed aswell as as-deposited films will be given in Chapter 5.

Many different bath compositions (different refers to bath constituentsrather than simply to different concentrations) have been used to deposit PbS andPbSe, more so than the II–VI materials. Both alkaline and acid baths have beendescribed, although the former are much more common. The chalcogen sourcesare similar to those used for the II–VI compounds, including thiourea, most com-monly used for PbS and selenourea (originally the source of choice) and, morecommonly nowadays, selenosulphate for PbSe. The variety of complexing agentsused is large; various carboxylic acids (most commonly citrate), triethanolamine,nitrilotriacetate, hydroxide, and even selenosulphate itself have all been used. Ad-dition of thiosulphate to a citrate/selenosulphate bath for PbSe resulted in a de-posit that did not show an XRD pattern and was assumed to be amorphous; blueshifts in the optical spectra could be explained either by structural changes or, pos-sibly, by size quantization in crystals too small to be seen in the XRD spectrum.

PbS has been deposited from an acidic bath using thiosulphate. Also, verythin films of PbS (maximum absorbance �0.015) have been grown on quartz im-mersed overnight in a solution of Pb ions together with polyvinyl alcohol (ostensi-bly to protect against aggregation of colloids) through which H2S has been passed.

2.9.2.2 Film Structure and Morphology

PbS and PbSe are almost always found in the rocksalt (RS) crystal form. All struc-tural investigations on CD films have shown this form, with one exception; PbSedeposited from hydroxide complex at high hydroxide concentrations and at rela-


tively high temperatures have produced a novel rhombohedral modification withan external hexagonal morphology and crystals typically 1–3 �m in size. ForPbSe, crystal size varies considerably, depending on deposition parameters, moreso than for the II–VI compounds. Crystals as small as 4 nm and larger than a mi-cron can be grown. The main factors, as for the II–VI materials, are temperatureand complex:Pb ratio. The use of a hydroxide complex gives the widest range ofcrystal sizes. PbSe often exhibits a bimodal distribution of crystal sizes. PbS, incontrast, tends to form only relatively large crystals.

Chemical deposition films of PbS and PbSe are generally not strongly tex-tured. One report has described (200) textured PbS films on glass if H2O2 is pre-sent in the deposition solution.

2.9.2.3 Electrical Properties

The IV–VI films are usually p-type, both as deposited and after annealing in air.One study, where PbS was deposited from a bath containing hydrazine, found thedeposit on glass to be n-type temporarily but converted to p-type on air exposure.By depositing the PbS on a trivalent metal coating (such as Al), the n-type con-ductivity could be stabilized for a longer time.

Electrical resistivity of the films, both PbS and PbSe, has often been re-ported to be of the order of 105 �-cm as deposited, with a reduction of about anorder of magnitude after annealing in air. However, this can vary considerablyfrom one type of deposition to another. Resistivities greater than 109 �-cm havebeen reported in some cases, which invariably drop to the k�-cm range after airannealing. These high-resistivity films are probably those with a very small crys-tal size (small meaning ca. 10 nm or less).

There are many studies on photoconductivity in these films, many of themearly ones and focused on annealed films (since air annealing is necessary foroptimal photoconductivity). The use of a chemical oxidant (which never seemsto be specified) gives much higher photosensitivity for as-deposited films thanfor films deposited without oxidant, although even here annealing is used to ob-tain maximum performance. Some studies on photoconductivity in as-depositedPbSe films have shown shifts in photoconductivity spectral response, with on-sets shifted to 2.2 �m instead of the ca. 4.5 �m more typical of annealed films.As with optical absorption studies, these shifts can be attributed to size quanti-zation.

2.9.3 Other Sulphides and Selenides

A large range of other metal sulphides and selenides have been deposited byCD. Since these will be individually described in Chapter 6, it will be sufficienthere to list all binary sulphides and selenides (along with oxides) in Table 2.1,along with up to three references to each compound.


2.9.4 Oxides, halides, and elemental Se

2.9.4.1 Oxides

While most studies on CD have been on sulphides and selenides, considerablework has also been carried out on oxide films. The films are most often formed byreaction of hydroxide ions with a metal salt. While it might be expected that theproduct is a hydroxide rather than an oxide, in many reported cases oxides are di-rectly formed. This is probably due to two factors: Many of the metal ions used(e.g., Pb, Sn, Tl,) do not readily, if at all, form simple hydroxides; others (Ag, Cu,Mn) are very readily oxidized even in aqueous solutions. Ni(II) hydroxide is fairlyreadily dehydrated, particularly in the presence of the persulphate ion used to de-posit the oxides in some cases. Zn(OH)2, Cd(OH)2 and In(OH)3 are reasonablystable; Zn(OH)2 can be easily dehydrated while the other two require annealing toform the oxide.

TABLE 2.1 Binary Semiconductors Deposited by CD

Sulphides Selenides Oxides

Ag2S [35–37] Ag2Se [38, 39] Ag2O/AgO [40, 41]As2S3 [42, 43]Bi2S3 [44–46] Bi2Se3 [47–49]CdS [50–52] CdSe [20, 53, 54] CdO [55–57]CoS [58] CoSe [59] CoO [60, 61]CuxS [62–64] CuxSe [65–67] Cu2O [68]

FeO(OH) [29]Fe2O3 [69]Fe3O4 [70]

HgS [71–73] HgSe [38, 74, 75]In2S3 [76–78] In2O3 [79, 80]MnS [81–83] MnO2 [84, 85]Mo-S [86,87] Mo-Se [86-88]NiS [89] NiSe [89] NiO [90, 91]PbS [92–94] PbSe [53, 95, 96] PbO2 [97, 98]Sb2S3 [43, 99, 100] Sb2Se3 [101–103] Sb2O3 [102]SnS [104–106] (also Sn2S3) SnSe [107]SnS2 [42, 106, 108] SnO2 [109–111]

TiO2 [32, 112, 113]TlS [114] TlSe [115,116] Tl2O3 [84, 98]ZnS [117–119] ZnSe [120–122] ZnO [123-125]

Only three references to pure CD deposition of tellurides, all of them for CdTe, have been found[126–128] and therefore no column for tellurides is given here. SiO2, Y2O3, and ZrO2, not given in thetable, have also been deposited, as have Ag halides.


Persulphates (also called peroxydisulphates) (S2O82�) are very strong oxi-

dizing agents and have been used to deposit oxides, in particular those withoxidation states higher than those of the original cation. PbO2, MnO2, and Tl2O3

deposited from solutions of Pb2�, Mn2�, and Tl� are examples of this type ofdeposition. Sometimes a small concentration of Ag� ions is needed in thedeposition solution as a catalyst. Ag� is a known catalyst for oxidationsusing S2O8

2�(S2O82� oxidizes it to Ag(III), and this is then the active oxidizing

agent).Fe-oxides, ZnO and In2O3 have been deposited using dimethylamineborane

(DMAB) or trimethylamineborane and Zn or In nitrate. The nitrate anion is im-portant and is believed to be reduced by the DMAB to nitrite and hydroxide. It isnot clear why ZnO, and not the hydroxide, is the final product. In the case of In,the deposited films were In(OH)3 and required annealing at ca. 200°C to form theoxide.

Homogeneous precipitation using urea (which hydrolyzes to give an alka-line solution) has been used extensively, and in a few cases films of basic salts(sulphates of Al and Sn(IV) and formate of Fe) have been obtained. These are notconsidered semiconductors in the conventional sense, but do provide examples forextension of the CD method beyond the conventional sulphide-selenide-oxidecompounds.

Many (hydr)oxides have been deposited by slowing down natural hydroly-sis, usually by complexation (e.g., AgO from an alkaline triethanolamine-com-plexed silver bath). Highly-acidic cations will readily hydrolyse even under acidicconditions. Fluoro-complexes of some of these (e.g., Ti, Si) can be controllablyhydrolyzed by addition of boric acid, which reacts with the F, thereby destroyingthe complex and allowing hydrolysis.

A (not-closed) heated solution of ammonia gradually loses ammonia. If acation is complexed with ammonia, the free-cation concentration will graduallyincrease as ammonia is lost (a rare example of slow cation release rather thananion release). It will also increase with an increase in temperature, due to thedecrease in stability of complexes with increasing temperature. As one exampleusing this principle, thin films of mixed ZnO/Zn(OH)2, which converted to ZnOon heating over 200°C, were deposited from a heated aqueous Zn-ion/ammoniabath.

Apart from these methods, there are others that are relevant to this section.Aqueous solutions of permanganate will slowly decompose, forming a brown film[MnO2 or possibly MnO(OH)] on the walls of the vessel in which they are stored.Increase of either acidity or alkalinity of the solution can accelerate this decom-position reaction. As another example of film formation due to slow solution de-composition, the author possesses a glass bottle with a green, highly tenaciousfilm of (presumably) Cr2O3 resulting from years of storage of some (unfortunatelyunknown) Cr-containing solution.


Since one common use of oxide films is for transparent, conducting coat-ings, the resistivities of these films were usually measured. Table 2.2 shows somebasic electrical and optical properties of some of these films.

2.9.4.2 Halides

Silver halides, in particular AgI and AgBr, have been deposited by hydrolysis ofhalogenoalcohols (halohydrins) to free halide ions in a solution of Ag� under

TABLE 2.2 Electrical and Optical Properties of CD Oxidesa

Resistivity Special Conductivity AbsorptionMaterial (�-cm) conditions type onset (nm) Reference

Cu2O 10�1 p ca. 600 68Fe3O4 2 � 103 Black 129Fe2O3 2 Annealed ca. 570 130

at 350�CIn2O3 109 As deposited n ca. 400 80

33 Annealedat 200�C

In2O3 2.10�2 79MnO2 pH � 6.3 84

pH � 8NiO 105 p ca. 700 90NiO 3 � 102 �430 131PbO2 10�3 1.7 98SnO2 10�1 Annealed n 350 109

at 250�C2 � 10�3 Annealed n

at 400�CSnO2 10�3 Annealed n 310

at 250�C(5% Sb) 10�4 Annealed n

at 400�CTiO2 109 32Tl2O3 4 � 10�4 600 98ZnO 2 � 104 to 4 � 102 Depending n �380 132

on boroncontent

ZnO �10�2 450 133ZnO:Al 2 � 10�4 to 2 � 10�2 340

In contrast to the rest of this book, optical properties are given as an approximate absorption edge in theabsorption (transmission) spectrum, rather than as a value of bandgap. This gives those who are notfamiliar with semiconductors a better feel for the appearance of the film.a As deposited, unless stated otherwise in the “Special conditions” column.


acidic conditions. For AgCl, a better method was found to be simple precipitationby adding a solution of NaCl to one of AgNO3. The concentration of the solutionsis important for this latter deposition—ca. 20 mM (much larger concentrations re-sult in rapid aggregation and little film formation, while much lower concentra-tions give a very slow deposition rate). A visible AgCl film (visible by scattering)is formed very rapidly—in seconds—which is unique in CD, which normally re-quires much longer to form a visible film. The crystal size of these silver halidefilms is large compared to most CD films, rarely less than 100 nm and sometimesas large as a micron.

2.9.4.3 Elemental Se

Virtually all the semiconductors deposited by CD are compound semiconduc-tors, the one exception being elemental Se. This has been deposited from solu-tions of selenosulphate, which rapidly form Se if acidified. By control of the pH,this reaction can be controlled to allow Se deposition to occur. Se films havealso been deposited from colloids of Se (prepared by reducing SeO2 solutions)by photodeposition, whereby the light activates the formation of films.

2.9.5 Coprecipitation of Metal Chalcogenides—Ternary Compounds

Chemical deposition is not limited to binary compounds. Ternary (and higher)compounds can be deposited by this technique. For the same reason as for the nonII–VI and IV–VI compounds in Section 2.9.3, this section will suffice with a tableof ternary compounds reported up to now, with two additions. The first is a briefconsideration of the principles involved in the deposition of materials containingthree or more elements. The second is to identify, in the table, which deposits havebeen clearly demonstrated to be a true single-phase solid solution rather than amixture of two or more phases.

If, e.g., thiourea is added to a mixture of Cd and Zn ions complexed withammonia, then, depending on the mechanism and experimental conditions of de-position, the deposit could be CdS, ZnS, a mixture of the two, a single-phaseCdxZnxS compound, or some combination of these. Structural characterization,most commonly XRD, together with elemental analysis, can usually reveal the na-ture of the product.

There a number of factors involved in determining just what the productturns out to be. An obvious one is consideration of the solubility of the variousproducts (and intermediates): The lower the value of Ksp of one binary product rel-ative to another, the more likely that product is (in principle) to deposit preferen-tially. This simple consideration is complicated by a number of other factors. Oneis the tendency of metal ions to coadsorb on the (usually high-surface-area) pri-


mary deposit, even if that metal chalcogenide (or whatever compound) is muchmore soluble than the primary one. Such adsorbed species can become occludedin the growing crystal as subsequent layers are built up. Another possibility, mostlikely for codeposition of structurally similar materials, is formation of a true solidsolution. Even if the relevant binary materials are not structurally similar, a truesingle-phase ternary deposit may be obtained, although in this case the composi-tion range is likely to be much narrower than in the previous case. For codeposi-tion of two different cations, compound formation may occur, particularly if thetwo cations have quite different acidic and basic properties. Chapter 8 describesthese principles in more detail.

Table 2.3 lists ternaries that have been deposited, together with indicationof when clear single compounds formation was verified. While solid solutionformation is usually the goal of these studies, it should be kept in mind that separate phases, either as a composite or as separate layers, may be required for some purposes. For example, bilayers of CdS/ZnO and CdS/ZnS have been deposited from single solutions. These depositions depend on the prefer-ential deposition of CdS over ZnS and, in the case of the former, the often-en-countered greater ease of formation of the oxide (hydroxide) than the sulphideof Zn.

TABLE 2.3 Ternary Materials Deposited by CD

Material References Material References

(Cd, Bi)S 134 In(S, OH) 135, 136*, 137, 138(Cd, Hg)S 139, 140* (Pb, Hg)S 144–144*(Cd, Pb)S 140, 145, 146–150* Pb(S, Se) 151*Cd(S, Se) 152–154*, 155, 156a (Sb, Bi)2S3 157* (after anneal)(Cd, Zn)S 158, 159*, 160, 161, Sn(S, OH) 165

162–164*(Cd, Zn)Se 166*, 167 (Zn, Cd)O 168*CuBiS2 169 Zn(O, OH), 170

Zn(S, O, OH)Cu3BiS3 171* (after anneal) Zn(S, OH) 170, 172CuInS2 173, 174*b Zn(Se, OH) 175CuInSe2 176–180* Zn(S, Se) 181, 182*Bi2 (S, Se)3 183* (Pb, Cu)S 184 (two phase)(Pb, Sn)Se 185* (Cd, Sn)O 133*(Pb, Bi)S 186*

Where a predominantly single phase (even over only part of the composition range) was at leastreasonably clearly demonstrated or could be inferred from the results, at least at some composition, therelevant reference is followed by an asterix, although sometimes this refers to annealed films.a Probably solid solution, based on Ref. 154.b See relevant section in Chapter 8.


2.10 NONAQUEOUS DEPOSITIONS

2.10.1 Introduction

Although almost all CD processes have been carried out in aqueous solution, thereare a few examples in nonaqueous media (specifically, in carboxylic acids). It isworth considering the differences between aqueous and nonaqueous deposition ingeneral first.

Solubility is an important criterion. The reagents involved need to be solu-ble in the solvent (this is, of course, obvious, but sometimes the obvious needspointing out). This limits the choice of reagents, compared to aqueous solutions.For example, sodium selenosulphate, the most common selenide-producingreagent, is insoluble in most organic solvents at any useful concentration. Sele-nourea (or one of its derivatives) is more useful in this case. When considering thechoice of metal salt, halides and particularly iodides as well as perchlorates tendto be more soluble in organic solvents than most other common anions.

Next there is the question of whether the anion-producing reaction, which isnormally a hydrolysis, can occur in the absence of water. Alternate reactions maybe needed. This is not a problem for depositions which occur by the complex de-composition mechanism.

Finally, even if these criteria are satisfied, there remains the question ofwhether the product will adhere to form a film or just precipitate homogeneouslyin the solution. This is the most difficult criterion to answer a priori. The hydrox-ide and/or oxy groups present on many substrates in aqueous solutions are likelyto be quite different in a nonaqueous solvent (depending on whether hydroxidegroups are present or not). Another factor that could conceivably explain the gen-eral lack of film formation in many organic solvents is the lower Hamaker con-stant of water compared with many other liquids; this means that the interactionbetween a particle in the solvent and a solid surface will be somewhat more in wa-ter than in most other liquids (see Chapter 1, van der Waals forces). From the au-thor’s own experience, although slow precipitation can be readily accomplishedfrom nonaqueous solutions, film formation appears to be the exception rather thanthe rule. The few examples described in the literature are confined to carboxylicacid solvents (see later).

What are the advantages of deposition from nonaqueous solutions? One isthe possibility to form films of compounds that are soluble, or not sufficiently in-soluble, in water. A (potential) example of this is formation of halides using, e.g.,chlorohydrins, which are in general soluble in organic solvents, to generate Cl�.

This could be expanded to materials that tend to form hydroxy-chalcogenidesrather than pure chalcogenides, such as ZnS and In2S3. If water (including water ofhydration from the salts) is rigorously excluded, the hydroxy impurity cannot formfrom solvents that do not contain hydroxy groups.


2.10.2 Examples of CD from NonaqueousSolutions

Films of various stoichiometries of Sn-S have been deposited from carboxylicacid (acetic, propionic, butyric) solutions of elemental S and SnCl2. Depending ondeposition conditions, in particular whether some water was present and howmuch, as well as the presence of complexing agents, films of approximate com-position SnS, Sn2S3, or SnS2 could be formed. Interestingly, various Sn-S filmswere also formed on the walls of the deposition vessel above the liquid level (byseveral centimeters); this was attributed to reaction between volatile SnCl4 andH2S, both formed in the deposition bath.

Bi2S3 was deposited from glacial acetic acid solutions of Bi(NO3)3 mixedwith formaldehyde solutions of Na2S2O3 or thioacetamide. Sb2S3 films were de-posited in a similar manner from acetic acid solutions.

These depositions are described in more detail in Chapter 6.

2.11 RELATED DEPOSITION TECHNIQUES

The CD technique is based on either slow formation of a reactive anion or slowdecomposition of a complex compound. However, there are other techniques, notinvolving these slow steps, that nevertheless are sometimes called chemical de-position or chemical solution deposition or are closely related. These techniquesare dealt with very briefly in this section.

2.11.1 Successive Ion-Layer Adsorption andReaction (SILAR) Process

The SILAR process is, as its name suggests, somewhat analogous to molecularbeam epitaxy (MBE), although the films obtained are most often not epitaxial.Like MBE, SILAR proceeds via a layer-by-layer buildup of the film, except in so-lution instead of in a vacuum. In the SILAR process, the substrate is immersedfirst in a solution containing the metal cation, rinsed, then immersed in a solutioncontaining the desired anion, and again rinsed. This gives (ideally) one monolayerof the deposit. The process is then repeated for as many times as needed to obtainthe required thickness. (It is probably not important if the anion or the cation is ad-sorbed first; after the first cycle, the process should be the same, although it is con-ceivable that differences in adhesion to the substrate may result, depending on theinitial order.) The rinsing steps are important, since without them relatively largereservoirs of one ion would remain on the substrate, and clusters of the semicon-ductor, rather than a film, would result. In fact, by omitting the rinsing step, filmscan be built up much more quickly. Thus, by successively immersing a glass slidein fairly strong solutions of Na2S and a Cd salt (say, 0.1M), visible yellow films


of CdS appear after only several cycles. However, such films are less likely to behomogeneous, compared with a properly prepared SILAR film. Cu2O, ZnO, andSnS have been deposited in this way (see Ref. 187 and references therein). Thetrue SILAR process is based on the expectation that, after each rinse step, only onemonolayer of the previously adsorbed ion will remain, which will favor a layer-type growth.

The technique is slow and tedious, but automation of the process can be car-ried out, whereby the substrate is attached to stepping motors that alternately im-merse and remove it from a series of beakers. This is shown schematically in Fig.2.7 for the example of CdS. The substrate is attached to an arm that can be movedboth vertically (to immerse/remove it) and horizontally (or in a circle) to positionit above different reaction vessels.

While the layers are usually polycrystalline, epitaxial layers of bothzincblende and wurtzite CdS have been grown on various single-crystal substrateswith lattice parameters close to those of CdS, as might be expected for an ion-by-ion growth.

Besides CdS, many other semiconductors have been deposited by theSILAR technique as well as organic conducting polymers, such as polypyrroleand polyaniline. For representative references and to locate some groups workingin this field, see Refs. 188–195.

2.11.2 Pyrolysis of Precursor Films

This method is based on pyrolysis of a metal chalcogenide–containing precur-sor. Heating CD hydroxide films to form oxides is a simple example and is com-mon in the deposition of some oxides. This will be treated in Chapter 7, which

FIG. 2.7 Simplified schematic diagram of an automated SILAR process for CdS.


deals with oxides. Somewhat further from CD as understood in this book, themethod is also widely used as the well-known sol-gel method for oxide films, inparticular SiO2 and TiO2. In the sol-gel method, a more-or-less viscous colloidalgel is deposited on a substrate either by dipping the substrate in the gel andslowly raising it from the solution or by spin-coating. The film, whose thicknesscan be controlled by the viscosity of the sol and either the rate of removal fromsolution or the rotational speed of the spin-coater, is then pyrolyzed to form theoxide.

This technique is not so easily extended to nonoxides. Sulphur, for example,does not form the cross-linking bonds needed to form the sol-gel as readily as doesoxygen. However, a related method has been used, albeit to a very small extent, toform CdS films. It is based on the thermal decomposition (at ca. 300°C) of aCd–thiourea complex, which is formed as a film by slowly withdrawing the sub-strate from a methanolic solution of a Cd salt and thiourea [196].

This can certainly be extended to other metal sulphides, using other com-plexes of sulphur (and also selenium). However, the complex and anion of themetal salt need to be chosen so that all the by-products of the pyrolysis reactionare volatile, otherwise the film will be contaminated with the nonvolatile by-prod-ucts. For example, using cadmium nitrate and thiourea, all the by-products arevolatile:

Cd(NO3)2 � (NH2)2CS → CdS � CO2 � 2H2O � 2N2O (2.20)

2.11.3 Exchange Reactions

Exchange reactions to convert one material into another by immersion in suitablesolutions are well known. Such a reaction is the basis for formation of the once-popular CdS/Cu2S photovoltaic cell, where a CdS film was immersed in a CuClsolution and part of the film converted to Cu2S. Since the solubility product ofCu2S is lower than that of CdS (or, in electrochemical language, Cu is more noblethan Cd), exchange is thermodynamically favorable. Examples of such exchangereactions using CD films are CdS and CdSe to the corresponding silver salts [197]and SnS2 to Ag2S [198].

2.11.4 Intermixing by Annealing Multilayers

By depositing two (or more) different layers and annealing them, intermixing ofthe layers can lead to ternary and multinary compounds, although clear compoundformation does not always occur. Thus, annealing (at 150°C, a relatively low tem-perature) ZnS-CuxS and PbS-CuxS films resulted in extensive interdiffusion of themetallic elements but no XRD confirmation of solid solution formation [199]. Onthe other hand, Sb2S3-CuS layers converted fully to CuSbS2 at 400°C, which ex-


hibited p-type conductivity and a direct bandgap of 1.5 eV [200]. By evaporatingIn onto CD Sb2S3, InSb was formed after annealing at 300°C.

2.11.5 Gas–solution Interface Reaction

Films of semiconductors up to a few hundred nanometers in thickness have beenformed by the simple reaction between H2S or H2Se and the surface of an aqueoussolution of the metal ion. This technique was described long ago for PbS and PbSe[201]. A more comprehensive description of the method, extended to a number ofdifferent metal sulphides, has been given [202]. It is stressed that the gas phase ispassed over the solution surface and not bubbled through the solution, which wouldbreak up the film. In fact, if the gas flow continues for too long (typically more thana few minutes), the film tends to break up and precipitate. Since the substrate forthese films is a liquid surface, the films can be (carefully) picked up and transferredto another surface or possibly even be self-supporting in small areas.

2.12 SPECIFIC TOPICS RELEVANT FOR CD FILMS

2.12.1 Solar Cells

2.12.1.1 Solid-State Photovoltaic Cells

Probably the most important factor responsible for the renewal of interest in theCD technique is the almost universal use of CD CdS films in thin-film photo-voltaic cells based on either Cu(Ga)InSe2 (abbreviated here as CIS, which in-

FIG. 2.8 Schematic diagram of the CIS/CdS and CdTe/CdS photovoltaic cells. The backcontact to the CdTe cell—Cu-doped carbon paste – is a commonly used one, but there areseveral modifications to this contact as well as completely different ones in use.


cludes Cu(In,Ga)Se2, usually abbreviated in the literature as CIGS, CuInSe2, andCuInS2) or CdTe absorber films. Schematic diagrams of these cells are shown inFigure 2.8. The common component of both is a thin (50–100 nm) film of CD CdScalled the window layer [since incident light (ideally) passes through it] or thebuffer layer (which points out the lack of understanding of exactly what this layerdoes). The CIS cell is a substrate (frontwall) cell (light passes through the semi-conductors in the direction of the substrate), while the CdTe cell is normally in thesuperstrate (backwall) configuration (light passes first through the substrate).

With only a few exceptions, the CdS is deposited from a standard ammo-nia/thiourea bath at ca. 70°C, with variations in the concentrations of reactants, theuse of temperature programming, and some variation in pH (using an ammoniumchloride buffer). It is notable that, in spite of many attempts to substitute CdS byanother CD material (driven by the desire for a environmentally friendlier mate-rial), CdS remains the best material to date for this purpose, both for CIS and CdTecells. Other materials deposited by CD include various Zn(OH,S), Zn(OH,Se),and In(OH,S) “compounds” and In(OH)3. The three first materials appear to be in-completely sulphided or selenided hydroxides, and it is not clear whether they area mixed or a single phase. Also, it is usually unclear whether oxide or hydroxidealso occurs [although one XPS study of In(OH,S) has demonstrated the absenceof either In(OH)3 or In2S3 in the film]. While some of these buffer layers approachCdS in terms of cell efficiency, they are invariably inferior.

Studies have been undertaken in an attempt to understand why CdS appearsto be so unique in this role and why CD is the best technique to deposit it. The CDsolution clearly plays an important role, not only in depositing the CdS, but alsoin its effect on the absorber surface. In the case of CuInSe2, the solution removesnative oxides (of In, Cu, and Se), removes excess CuxSe, and forms an interfacethat is Cu-deficient and contains Cd. In fact, Cd has been shown to diffuse a smallbut appreciable distance (ca. 10 nm) into Cu(Ga)InSe2 films but not into CuInSe2

single crystals. It is not clear whether this is due to differences in composition orin crystallinity. However, the diffusion of Cd into CIS is believed to be related tothe presence of a Cu-depleted region at the CIS surface; the CD bath, as alreadynoted, is instrumental in forming such a region. Additionally, it is thought that re-moval of surface oxygen substituting for Se vacancies by replacement with S mayincrease band bending by modification of the surface charge.

Since the envisionaged application of a CD process in thin-film solar cellsis a large-scale one, efforts have been made to optimize the deposition processused, particularly in minimizing the waste Cd-containing solutions. Dilute Cd so-lutions (ca. 1 mM), a flow system with filtration, and a heated substrate have beenemployed to this end. The heated substrate means that deposition occurs prefer-entially on the substrate rather than on the cooler walls of the deposition vessel.Also, ethylenediamine has been used as a complexant rather than the much morevolatile ammonia.


There are a number of studies on other photovoltaic properties of CD films(see Chapter 9). As an example, p-n junctions have been fabricated by depositingPbS on a glass substrate partially coated with a trivalent metal, such as Al. ThePbS deposited on the glass is (as usual) p-type, while that deposited on the metalcoating was n-type, at least for some time. Photovoltages up to 0.1 V (at room tem-perature) and 0.28 V (at 90 K) were measured (the bandgap of PbS is ca. 0.4 eV –even lower at low temperatures so only small photovoltages are expected).

2.12.1.2 Photoelectrochemical Cells

One of the attractive features of CD is its simplicity (in terms of carrying out thedeposition, that is, not always in understanding the deposition itself). The sameproperty of simplicity is often ascribed to photoelectrochemical cells (PECs).Therefore it is not surprising that CD has often been used to fabricate the semi-conductor electrodes for PECs.

A PEC is the liquid junction analogue of a solid-state Schottky cell. In itssimplest configuration, a semiconductor with an ohmic contact is immersed in anelectrolyte and connected via a load to a second counterelectrode (often platinumor graphite). Super-bandgap light incident on the semiconductor creates elec-tron/hole pairs that are separated by the electric field (space charge layer) formedby the contact between the semiconductor and the electrolyte. One charge type(holes, for the more common case of n-type semiconductors) is injected into theelectrolyte. The holes oxidize some electrolyte species, while the electrons are ex-tracted through the ohmic contact and flow through an external circuit to the coun-terelectrode, where a reduction occurs. If the same species is both oxidized (at thesemiconductor electrode) and reduced (at the counterelectrode), no net change oc-curs in the electrolyte, and electrical energy is generated in the external circuit.This type of cell is called a regenerative PEC. If different species are electrolyzed,chemical energy can be produced. Most PECs made using CD semiconductorfilms are of the former type.

In nanocrystalline semiconductor films (commonly obtained in CD), thecrystal size may be too small to support an appreciable space charge layer.Charges in that case are separated by differing kinetics between electron and holeinjection into the electrolyte. The upcoming discussion on nonannealed filmstreats this in somewhat more detail. (Chapter 9 discusses PECs and their princi-ples of operation more fully.)

The majority of PEC studies have been carried out on either CdS or CdSefilms, although many other CD semiconductors have been demonstrated to exhibitPEC activity. In most, though not all, cases, these films were annealed for opti-mum PEC response. Films annealed at temperatures above ca. 300°C usually ex-hibit a large degree of crystal growth, and therefore such films will be discussedseparately from as-deposited films, which, in most cases, are composed of crys-tals �20 nm in size.


Annealed Films. For many years, the CdSe photoelectrode—polysulphideelectrolyte PEC—was probably the most studied system in PEC research. Thebandgap of (bulk—see later) CdSe is ca. 1.7 eV, which is close to the theoreticaloptimum of 1.5 eV for photovoltaic cells in general; and in relative terms, that sys-tem was fairly stable in terms of self-oxidation of the semiconductor film in theelectrolyte by the photogenerated holes.

Simulated solar conversion efficiencies up to 6.8% on Ti substrates havebeen reported for annealed CD CdSe films in polysulphide electrolyte based on alow-ammonia-concentration–selenosulphate bath. Several successive depositionswere required to build up an optimum final film thickness of 2.5 �m (when mostof the light was absorbed). The initial deposit was annealed to improve adherenceand the final multideposited film was annealed at 550°C in air, followed by etch-ing and zinc ion treatment.

By incorporating silicotungstic acid (STA) in the deposition bath, largerparticle size (after annealing at 430°C) and much better conversion efficiencies(compared to the STA-free bath) were obtained—11.7% based on tungsten-halo-gen illumination (solar efficiency is lower). The effect of the improved perfor-mance was not clear. It was suggested that the STA improved charge transfer ki-netics at the CdSe/electrolyte interface. The larger particle size may also be animportant factor.

Chemical deposition CD CdS has shown much lower efficiencies in a PEC.This is due to its higher bandgap, which allows only a small fraction of solar ra-diation to be absorbed in the film.

Nonannealed films. Although the conversion efficiencies are much lowerthan those of annealed films, the PEC properties of as-deposited films show otherinteresting properties, connected with their nanocrystalline and somewhat porousmorphology. As already noted, there is usually no appreciable space charge layerin these nanocrystals. Since the films are porous and electrolyte can reach (mostof) the surface of all the nanocrystals, charge generation can be considered to takeplace at the surface of the individual crystals—there is no need for a field to pro-vide the driving force for charge drift. Electrons and holes are then separated bydifferent kinetics for electron and hole transfer to the electrolyte, which in turn isaffected by the relative trapping efficiencies of the charges at the nanocrystal sur-face. If one charge (say, holes) is removed rapidly by the electrolyte, the electronscan get to the back contact without recombination. In practice, much of the pho-togenerated charge is lost, probably by injection of both carriers into the elec-trolyte (indirect recombination).

A unique property of these films results from this mechanism. CdSe, as de-posited, behaves as an n-type semiconductor (holes are transferred to the elec-trolyte, while electrons are extracted at the back contact). In this case, the holesare preferentially trapped at the surface, and are more readily injected into solu-


tion. Mild etching in dilute HCl changes the distribution of trapping states at thesurface in such a way that electrons are now more readily trapped and are prefer-entially injected into the (identical polysulphide) electrolyte, the film behaving asa p-type semiconductor. It is notable that this behavior, typical for a crystal size of4–5 nm, is not observed when the crystal size is 20 nm. The very high surface areaand high density of trapping states appear to be determining here.

As-deposited CdS has been studied as a photoelectrode with variousdopants (Al, As, Cu) incorporated in the deposition bath. The emphasis in thesestudies was on PEC efficiencies, which were very low in all cases, although dop-ing, particularly with As, did have a beneficial effect.

An interesting variation of photoactivity has been observed during the de-position of several semiconductor films (CdSe, CdS, PbS, Bi2S3). Illuminationduring deposition increases the deposition rate and, in some cases, increases thecrystal size somewhat. While some effects can be ascribed to local heating, in-creasing deposition rate, the main effect is probably electron-hole generation bysuper-bandgap light and reduction of the chalcogen species to chalcogenide by thephotogenerated electrons at the growing crystal surface, resulting in the formationof more metal chalcogenide.

Some other semiconductors have been investigated as photoelectrodes, inall cases giving low but appreciable photoactivity. These include Bi2S3, PbSe,Sb2Se3, SnS2, HgS, and Ag2S. Chemical deposition has also been used to formboth the photoelectrode (CdSe) and the storage electrode (Ag2S) in a photoelec-trochemical storage cell, where the Ag2S acts as an in situ storage electrode.

2.12.2 Quantum Size Effects in CD Films

Films of materials deposited at or near room temperature (and in this respect100°C is considered to be near room temperature) tend to have a small crystal size.This is not surprising since high temperatures are normally required to impart suf-ficient mobility to a freshly deposited species in order for recrystallization to oc-cur. This small crystal size, which at one time was almost universally consideredto be a disadvantage, is increasingly considered to be an advantage as interest innanocrystalline and nanoparticle materials grows. The term nanocrystalline usu-ally refers to materials with a crystal size from a nanometer up to hundreds ofnanometers (at this upper limit, the term microcrystalline starts to take over).

The size of the crystals formed in CD films is often small enough that quan-tum-size effects become apparent. The terms quantum size effect and size quanti-zation are normally used to describe a material whose energy structure differsfrom that of the bulk material. As crystal (or, more generally, particle) size de-creases, charges (electrons and holes) in the particles are constrained in an in-creasingly small volume. When the particle size becomes smaller than the Bohrdiameter of the charges in the bulk material (between 2 and 20 nm for many ma-


terials), the quantum-size effect is manifested as an increase in the bandgap, Eg,of the material and separation of the energy bands into discrete levels. The in-crease in Eg is most commonly seen as a blue shift (i.e., to higher energy) of theoptical absorption spectrum (see Chapter 10 for more details on this topic.)

This effect is shown in Figure 2.9 for CdSe films deposited from baths con-taining Cd complexed with NTA (nitrilotriacetate) and Na2SeSO3 as a Se source.The nanocrystal size, measured by both XRD and TEM, varied from ca. 3 nm upto 20 nm with increase in temperature and/or change in mechanism from a clustermechanism to an ion-by-ion deposition. The optical bandgap shifts from 1.8 eV(for bulk, zincblende CdSe) to ca. 2.4 eV for the smallest nanocrystals (ca. 3 nm).

The main difference between the two mechanisms as they relate to crystalsize (discussed in Sec. 2.6) is that the cluster mechanism is three dimensionalwhile the ion-by-ion one is mainly two dimensional. Crystal size in the former islimited largely by the amount of reactant per nucleus: The more nuclei, the smallerthe final crystal size, since the same concentration of reactants is divided overmore nuclei. Temperature affects this by stabilizing (kinetically) smaller nuclei astemperature is lowered, thus increasing the number of nuclei at lower temperature,

FIG. 2.9 Transmission spectra of CD CdSe films deposited at various temperatures fromCdSO4/NTA/Na2SeSO3 baths. All samples deposited by hydroxide cluster mechanism ex-cept 80°C HC (high complex), which proceeded via the ion-by-ion mechanism. The effectivebandgap can be approximated by the wavelength (photon energy) a little shorter (higher) thanthe absorption onset. A second absorption knee, ca. 0.4 eV to higher energy of the initial on-set, seen clearly in the 41°C and 80°C samples, is due to a transition from the spin-orbit va-lence level to the lowest conduction level and is commonly observed in these films.


resulting in a smaller final crystal size. For the two-dimensional ion-by-ionmechanism, the final crystal size is probably determined by the number of nucleion the substrate (initially bare, eventually covered with deposit), which growuntil they touch other crystals and cannot grow further in the horizontal (to thesubstrate) direction. Increased temperature also may allow increased coalescenceof the individual crystals into larger ones.

Other CD semiconductors have been shown to exhibit size quantization.PbSe shows the effect very clearly, since quantum size effects can be clearly seenin this material, even in crystals up to several tens of nanometers in size (due tothe small effective mass of the excited electron-hole pair). Shifts of greater than 1eV have been demonstrated, from the bulk bandgap of 0.28 eV to 1.5 eV.

CdS, deposited from the usual ammonia/thiourea bath, normally gives filmsthat are not quantized, with crystal size larger than 10 nm. Some exceptions do oc-cur, however. Small increases in bandgap have been found when high thioureaconcentrations were used. At high values of pH (�12), smaller crystal sizes (downto ca. 5 nm) have been obtained with blue spectral shifts corresponding to in-creases in bandgap of up to a few hundred meV for very thin films. The crystalsize increased with increase in film thickness. Using NTA as a complex and work-ing under conditions where the cluster mechanism is operative, 5-nm nanocrystalsof CdS exhibiting quantum size effects have also been obtained. This crystal sizedid not vary much with film thickness. Using a process of electrochemically in-duced CD of CdS (see Chapter 4), nanocrystalline CdS films were grown using 2-mercaptoethanol as a strongly adsorbing growth-termination (capping) agent. Byincreasing the concentration of mercaptoethanol, crystal size was reduced downto 4.1 nm (and bandgap increased up to 2.69 eV).

Both ZnS and ZnSe films have been grown that show moderate increases inbandgap (up to a few hundred meV). The Zn chalcogenides generally exhibitsmaller increases in bandgap than the corresponding Cd compounds of the samesize, due to their larger effective masses.

HgS, deposited from thiosulphate solution, has been described with crystalsizes that depend on deposition temperature, from 3 nm to 8 nm and correspond-ing variation in apparent bandgaps from 2.4 to 1.9 eV.

HgSe has been deposited by CD exhibiting different bandgaps and onesstrongly shifted from the bulk value (bulk HgSe is a semimetal with a negativebandgap). Values as high as 2.5 eV for 7.7 nm crystal size have been reported. An-other bath composition gave a bandgap of 1.42 eV, although this was not ex-plained through size quantization but because of an amorphous structure.

Bi2S3 has been deposited with film thickness–dependent crystal sizes, from5.2 to 8 nm and corresponding bandgaps, measured from the absorption spectra,from nearly 2.3 down to 1.8 eV.

Cu-S films, of various stoichiometries, have also shown small quantum sizeeffects.


Apart from these reported quantum size effects, there are a number of ma-terials that, while quantum size effects have not been explicitly invoked, havebeen reported to show blue spectral shifts that may be explained in the same man-ner. These include Cu1.8S, various CuS compositions (Eg 1.8—2 eV), Sb2S3, andSnS. These systems are all described in Chapter 10.

While shifts due to size quantization have most commonly been seen in ab-sorption spectroscopy, other spectroscopies, such as photoelectrochemical pho-tocurrent, photovoltage (using a vibrating Kelvin probe), photoluminescence, andphotoconductivity spectroscopies have all shown quantum shifts in various CDfilms.

2.13 APPLICATIONS (ACTUAL AND POTENTIAL)OF CD FILMS

We conclude this chapter with applications of CD semiconductor films, both thosethat have been realized and potential uses.

The most important use of CD films for many years was to make PbS andPbSe films for photoconductive detectors [10]. Such detectors, made by CD, arestill in use today, although they are facing competition from photovoltaic III–Vdetectors. It should be noted that for good photosensitivity, air-annealing of theCD films is carried out, and this annealing treatment is connected with partial oxidation of the PbS and PbSe surfaces.

Today, the the most important “application” for CD films is the use of CDCdS as the window (or buffer) layer in thin-film photovoltaic cells [16]. BothCdTe- and CuInSe2-type absorber films use this procedure. Such cells havereached the pilot plant stage, and there appears to be no obvious competitor for theCD CdS at present.

The study of CD semiconductors, and in particular CdSe and CdS, for useas photoelectrodes in photoelectrochemical cells is connected with this use, al-though much farther from likely application. This study was driven, to a large ex-tent, by the simplicity of deposition of the films, a particularly sought-after re-quirement for this purpose, both from the point of view of applications andbecause it allowed many groups (usually chemists) who maybe did not have readyaccess to conventional vacuum deposition systems to prepare the films.

The success of CD CdS in photovoltaic cells has driven related researchwith potential applications in other semiconductor devices. Since the CD processseems to play a role in the favorable properties of the CdS windows by decreas-ing interface recombination, studies of its passivation properties on other inter-faces and surfaces have been carried out, with considerable success. For example,when a very thin film (ca. 6 nm) was deposited between InP and SiO2, the result-ing reduction of the interface state density led to improved electrical properties ofmetal-insulator-semiconductor capacitors and field effect transistors (FETs)


based on this interface [203,204]. These improvements were correlated with re-moval of native oxides and protection against oxide regrowth by the CdS, as wellas with removal of phosphorus vacancies at the interface. Improvements were alsoobtained in InAlAs/InGaAs transistors and metal-semiconductor-metal photode-tectors by the same treatment [205]. It is clear that these effects are connected notonly with the properties of the CD films but also with reaction between the sub-strate semiconductor surface and the CD solution.

Another example of CD films applied toward electronic devices is the fab-rication of thin-film field effect transistors by depositing CdS (50 nm) onto oxi-dized n�-Si and annealing at 400°C [206].

Another potential application is for solar control coatings. Thin films of cer-tain sulphides—in particular those of copper and lead—are reasonably transmis-sive to visible solar radiation while reflecting most of the infrared radiation. Ifused as a window coating, the heat rejection from the solar radiation will result inreduced heating of the interior by solar radiation, compared to a noncoated glasswindow [23,24]. Such coatings are commonly in use, particularly in large build-ings, and often take on an attractive golden appearance by reflected light. They areapplied by vacuum techniques. If the material usage is sufficiently high (as hasbeen demonstrated for CD CdS used for photovoltaic cells), CD is an attractive al-ternative to deposit these coatings.

The high infrared reflectivity of these films implies highly doped (thereforerelatively highly electrically conducting) films. However, another type of solarcontrol coating, a solar shield for passive cooling, requires films that have a hightransmission in the mid-infrared. Passive cooling occurs when a surface emitsmore radiation than it absorbs [207]. The (cloudless) atmosphere is relativelytransparent to mid-infrared radiation (between ca. 8–13 �m—the atmosphericwindow); therefore, radiation of these wavelengths can be emitted from a surfaceat ambient temperature through the atmosphere to space, which is a sink at a verylow temperature (theoretically 4 K, in practice higher, but still very much lowerthan the surface of the Earth). This leads to a cooling effect at night and explainswhy car surfaces and grass become covered with much more dew during a clearnight than a cloudy one. However, during the day, solar radiation would swampthis cooling effect. To prevent this, a shield is required that is transparent to the8–13-�m region but that blocks the complete solar spectrum. Low-bandgap semi-conductors, such as PbS and PbSe on polyethylene (one of the very few materials,readily available in large areas, that is transparent to the atmospheric window re-gion), are suitable for this purpose [208], but they must be close to intrinsic to pre-vent free carrier absorption and reflection in the infrared region. Nanocrystallinefilms (and CD films are often nanocrystalline) are more likely to be intrinsic thanlarge-grain ones: unlike a macroscopic crystal, it is possible to obtain smallnanocrystals with zero doping concentration, i.e., truly intrinsic.

Other potential applications for CD films have been suggested and studied.


Related to the solar control coatings just described, solar absorbing coatings canbe used, e.g., for water heating. Multilayered stacks of CD PbS and CdS on Ni-coated Cu can be configured to minimize reflectivity in the solar region and min-imize emittance in the thermal infrared region [209]. The relatively high electri-cal conductivity of CD CuxS films has been exploited to form an electrical contactto ferroelectric films, as a partially transparent conducting film on plastic and alsoas a cupric ion sensor [210].

Our brief overview has given uses for CD films and rationales for theirstudy. It should be clear that more uses are likely to develop as the present resur-gence in interest increases the pool of knowledge in the field and allows deposi-tion of better-quality films, higher reproducibility, new materials and old ones indifferent forms than usual.

REFERENCES1. J Liebig. Ann. Pharmaz. 14:134, 1835.2. C Puscher. Dingl. J. 190:421, 1869.3. E Beutel, Z Angew. Chem. 26:700, 1913.4. E Beutel, A Kutzelnigg. Z. Elektrochem. 36:523, 1930.5. E Beutel, A Kutzelnigg. Monats. 58:295, 1931.6. J Emerson-Reynolds. J. Chem. Soc. 45:162, 1884.7. G Rosenheim, W Stadler, VJ Mayer. Z. Anorg. Chem. 49:1, 1906.8. G Rosenheim, W Stadler, VJ Mayer. Z. Anorg. Chem. 49:13, 1906.9. RH Bube. Photoconductivity of Solids. New York: Wiley, 1960.

10. DE Bode. In: Physics of Thin Films, Vol. 3. Academic Press, New York and Lon-don, 1966, p 275.

11. SG Mokrushin, YV Tkachev, Kolloidn Zh. 23:438, 1961.11a. ET Allen, JL Crenshaw, HE Merwin. Am. J. Sci. 34:341, 191212. H Uda, H Taniguchi, M Yoshida, T Yamashita. Jpn. J. Appl. Phys. 17:585, 1978.13. RW Birkmire, BE McCandless, WN Shafarman, RD Varrin Jr. In: 9th ECPV Solar

Energy Conf., Freiberg Germany, 1989, p 134.14. RH Mauch, M Ruckh, J Hedström, D Lincot, J Kessler, R Klinger, L Stolt, J Vedel,

H-W Schock. In: 10th ECPV Solar Energy Conf. Lisbon Portugal, 1991, p 1415.15. TL Chu, SS Chu, C Ferekides, CQ Wu, J Britt, C Wang J. Appl. Phys. 70:7608,

1991.16. TL Chu, SS Chu, N Schultz, C Wang, CQ Wu. J. Electrochem. Soc. 139:2443, 1992.17. G Hodes, A Albu-Yaron, F Decker, P Motisuke. Phys. Rev. B 36:4215, 1987.18. G Hodes, A Albu-Yaron. Proc. Electrochem. Soc., 88–14: 298, 1988.19. S Gorer, A Albu-Yaron, G Hodes. J. Phys. Chem. 99:16442, 1995.20. S Gorer, G Hodes. J. Phys. Chem. 98:5338, 1994.21. PK Nair, MTS Nair. J. Phys. D: Appl. Phys. 23:150, 1990.22. PK Nair, M Ocampo, A Fernandez, MTS Nair. Sol. Energy Mater. 20:235, 1990.23. PK Nair, MTS Nair, A Fernandez, M Ocampo. J. Phys. D: Appl. Phys. 22:829, 1989.24. PK Nair, MTS Nair, VM Garcia, OL Arenas, Y Pena, A Castillo, IT Ayala, O


Gomezdaza, A Sanchez, J Campos, H Hu, R Suarez, ME Rincon. Sol. Energy Mater.Sol. Cells 52:313, 1998.

25. HL Smith. J. Sci. Instrum. 4:115, 1927.26. AB Lundin, GA Kitaev. Inorg. Mater. 1:1900, 1965.27. P Pramanik, S Bhattacharya. J. Mater. Sci. Lett. 6:1105, 1987.28. FC Meldrum, J Flath, W Knoll. J. Mater. Chem. 9:711, 1999.29. PC Rieke, BD Marsh, LL Wood, BJ Tarasevich, J Liu, L Song, GE Fryxell. Lang-

muir 11:318, 1995.30. JH Fendler. Isr. J. Chem. 33:41, 1993.31. ML Breen, JT Woodward, DK Schwartz, AW Apblett. Chem. Mat. 10:710, 1998.32. RJ Collins, H Shin, MR DeGuire, AH Heuer, CN Sukenik. Appl. Phys. Lett. 69:860,

1996.33. XK Zhao, J Yang, LD Mccormick, JH Fendler. J. Phys. Chem. 96:9933, 1992.34. I Pintilie, E Pentia, L Pintilie, D Petre, C Constantin, T Botila. J. Appl. Phys.

78:1713, 1995.35. H Meherzi-Maghraoui, M Dachraoui, S Belgacem, KD Buhre, R Kunst, P Cowache,

D Lincot. Thin Solid Films 288:217, 1996.36. M Ristova, P Toshev. Thin Solid Films 216:274, 1992.37. SS Dhumure, CD Lokhande. Thin Solid Films 240:1, 1994.38. AA Velykanov, EK Ostrovskaya, NP Garina, VA Turacova, AA Tchurkan. Ukr.

Chim. Zh. 49:764, 1983.39. B Pejova, M Najdoski, I Grozdanov, SK Dey. Mater. Lett. 43:269, 2000.40. AJ Varkey, AF Fort. Sol. Energy Mater. Sol. Cell 29:253, 1993.41. T Kocareva, I Grozdanov, B Pejova. Mater. Lett. 47:319, 2001.42. CD Lokhande. Mater. Chem. Phys. 28:145, 1991.43. JD Desai, CD Lokhande. Thin Solid Films 249:135, 1994.44. S Biswas, A Mondal, D Mukherjee, P Pramanik. J. Electrochem. Soc. 133:48, 1986.45. MTS Nair, PK Nair. Semicond. Sci. Tech. 5:1225, 1990.46. CD Lokhande, AU Ubale, PS Patil. Thin Solid Films 302:1, 1997.47. P Pramanik, RN Bhattacharya, A Mondal. J. Electrochem. Soc. 127:1857, 1980.48. RN Bhattacharya, P Pramanik. J. Electrochem. Soc. 129:332, 1982.49. VM García, MTS Nair, PK Nair, RA Zingaro. Semicond. Sci. Technol. 12:645,

1997.50. PC Rieke, SB Bentjen. Chem. Mat. 5:43, 1993.51. R Ortega-Borges, D Lincot. J. Electrochem. Soc. 140:3464, 1993.52. JM Doña, J Herrero. J. Electrochem. Soc. 144:4081, 1997.53. RC Kainthla, DK Pandya, KL Chopra. J. Electrochem. Soc. 127:277, 1980.54. H Cachet, H Essaaidi, M Froment, G Maurin. J. Electroanal. Chem. 396:175, 1995.55. RL Call, NK Jaber, K Seshan, JR Whyte Jr. Sol. Energy Mater. 2:373, 1980.56. M Ocampo, AM Fernandez, PJ Sebastian. Semicond. Sci. Technol. 8:750, 1993.57. M Najdoski, I Grozdanov, B Minceva. J. Mater. Chem. 6:761, 1996.58. PK Basu, P Pramanik. J. Mater. Sci. Lett. 5:1216, 1986.59. P Pramanik, S Bhattacharya, PK Basu. Thin Solid Films 149:L81, 1987.60. AJ Varkey, AF Fort. Sol. Energy Mater. Sol. Cells 31:277, 1993.61. B Pejova, A Isahi, M Najdoski, I Grozdanov. Mater. Res. Bull. 36:161, 2001.62. I Grozdanov, M Najdoski. J. Solid State Chem. 114:469, 1995.


63. MTS Nair, L Guerrero, PK Nair. Semicond. Sci. Technol. 13:1164, 1998.64. P Pramanik, MA Akhter, PK Basu. J. Mater. Sci. Lett. 6:1277, 1987.65. A Mondal, P Pramanik. J. Solid State Chem. 47:81, 1983.66. I Grozdanov. Synth. Metals 63:213, 1994.67. C Lévy-Clément, M Neumann-Spallart, SK Haram, KSV Santhanam. Thin Solid

Films 302:12, 1997.68. I Grozdanov. Mater. Lett. 19:281, 1994.69. QW Chen, YT Qian, H Qian, ZY Chen, WB Wu, YH Zhang. Mater. Res. Bull.

30:443, 1995.70. QW Chen, YT Qian, ZY Chen, Y Xie, GE Zhou, YH Zhang. Mater. Lett. 24:85,

1995.71. KM Gadave, PP Kumbhar, CD Lokhande. Indian J. Pure Appl. Phys. 30:299, 1992.72. M Najdoski, I Grozdanov, SK Dey, B Siracevska. J. Mater. Chem. 8:2213, 1998.73. SS Kale, CD Lokhande. Mater. Chem. Phys. 59:242, 1999.74. P Pramanik, S Bhattacharya. Mater. Res. Bull. 24:945, 1989.75. B Pejova, M Najdoski, I Grozdanov, SK Dey. J. Mater. Chem. 11:2889, 1999.76. GA Kitaev, VI Dvoinin, AV Ust’yantseva, MN Belyaeva, LG Skornyakov. Inorg.

Mater. 12:1448, 1976.77. T Yoshida, K Yamaguchi, H Toyoda, K Akao, T Sugiura, H Minoura. Proc. Elec-

trochem. Soc. 97–20:37, 1997.78. R Bayón, C Guillén, MA Martínez, MT Gutiérrez, J Herrero. J. Electrochem. Soc.

145:2775, 1998.79. RP Goyal, D Raviendra, BRK Gupta. Phys. Status Solidi (a) 87:79, 1985.80. M Izaki. Electrochem. Solid State Lett. 1:215, 1998.81. P Pramanik, MA Akhter, PK Basu. Thin Solid Films 158:271, 1988.82. CD Lokhande, A Ennaoui, PS Patil, M Giersig, M Muller, K Diesner, H Tributsch.

Thin Solid Films 330:70, 1998.83. CD Lokhande, KM Gadave. Turkish J. Phys. 18:83, 1994.84. W Mindt. J. Electrochem. Soc. 118:93, 1971.85. KL Chopra, RC Kainthla, DK Pandya, AP Thakoor. In: G Hoss, MH Francombe,

JL Vossen, eds. Physics of Thin Films, Vol. 12. Academic Press, New York andLondon 1982, p. 167.

86. P Pramanik, RN Bhattacharya. J. Mater. Sci. Lett. 8:781, 1989.87. P Pramanik, S Bhattacharya. Mat. Res. Bull. 25:15, 1990.88. KC Mandal, O Savadogo. J. Mater. Chem. 1:301, 1991.89. P Pramanik, S Biswas. J. Solid State Chem. 65:145, 1986.90. P Pramanik, S Bhattacharya. J. Electrochem. Soc. 137:3869, 1990.91. AJ Varkey, AF Fort. Thin Solid Films 235:47, 1993.92. F Kicinski. Chem. Ind. 54, 1948.93. M Isshiki, T Endo, K Masumoto, Y Usui. J. Electrochem. Soc. 137:2697, 1990.94. KM Gadave, SA Jodgudri, CD Lokhande. Thin Solid Films 245:7, 1994.95. S Gorer, A Albu-Yaron, G Hodes. Chem. Mater. 7:1243, 1995.96. I Grozdanov, M Najdoski, SK Dey. Mater. Lett. 38:28, 1999.97. W Mindt. J. Electrochem. Soc. 117:615, 1970.98. RN Bhattacharya, P Pramanik. Bull. Mater. Sci. 2:287, 1980.99. KC Mandal, A Mondal. J. Phys. Chem. Solids 51:1339, 1990.


100. MTS Nair, Y Pena, J Campos, VM García, PK Nair. J. Electrochem. Soc. 145:2113,1998.

101. P Pramanik, RN Bhattacharya. J. Solid State Chem. 44:425, 1982.102. RN Bhattacharya, P Pramanik. J. Electrochem. Soc. 129:1642, 1982.103. RN Bhattacharya, P Pramanik. Sol. Energy Mater. 6:317, 1982.104. P Pramanik, PK Basu, S Biswas. Thin Solid Films 150:269, 1987.105. MTS Nair, PK Nair. Semicond. Sci. Tech. 6:132, 1991.106. RD Engelken, HE McCloud, C Lee, M Slayton, H Ghoreishi. J. Electrochem. Soc.

134:2696, 1987.107. P Pramanik, RN Bhattacharya. J. Mater. Sci. Lett. 7:1305, 1988.108. CD Lokhande. J. Phys. D: Appl. Phys. 23:1703, 1990.109. D Raviendra, JK Sharma. J. Phys. Chem. Solids 46:945, 1985.110. K Tsukuma, T Akiyama, H Imai. J. Non-Cryst. Sol. 210:48, 1997.111. QW Chen, YT Qian, ZY Chen, GE Zhou, YH Zhang. Thin Solid Films 264:25,

1995.112. QW Chen, YT Qian, ZY Chen, WB Wu, ZW Chen, GE Zhou, YH Zhang. Appl.

Phys. Lett. 66:1608, 1995.113. H Shin, RJ Collins, MR De Guire, AH Heuer, CN Sukenik. J. Mater. Res. 10:692,

1995.114. A Mondal, P Pramanik. Thin Solid Films 110:65, 1983.115. MJ Mangalam, KN Rao, N Rangarajan, CV Suryanarayana. Jpn. J. Appl. Phys.

8:1258, 1969.116. RN Bhattacharya, P Pramanik. Bull. Mater. Sci. 3:403, 1981.117. R Ortega Borges, D Lincot, J Vedel. In: 11th ECPV Solar Energy Conf. Montreux,

Switzerland, 1992, p 862.118. JM Doña, J Herrero. J. Electrochem. Soc. 141:205, 1994.119. P O’Brien, J McAleese. J. Mater. Chem. 8:2309, 1998.120. CD Lokhande, PS Patil, H Tributsch, A Ennaoui. Sol. Energy Mater. Sol. Cells

55:379, 1998.121. JM Doña, J Herrero. J. Electrochem. Soc. 142:764, 1995.122. P Pramanik, S Biswas. J. Electrochem. Soc. 133:350, 1986.123. P O’Brien, T Saeed, J Knowles. J. Mater. Chem. 6:1135, 1996.124. M Izaki, J Katayama. J. Electrochem. Soc. 147:210, 2000.125. A Ennaoui, M Weber, R Scheer, HJ Lewerenz. Sol. Energy Mater. Sol. Cells

54:277, 1998.126. GK Padam, SK Gupta. Appl. Phys. Lett. 53:865, 1988.127. RW Buckley. In: 11th ECPV Solar Energy Conf. Montreux, Switzerland, 1992, p

962.128. VB Patil, PD More, DS Sutrave, GS Shahane, RN Mulik, LP Deshmukh. Mater.

Chem. Phys. 65:282, 2000.129. M Izaki, O Shinoura. Adv. Mater. 13:142, 2001.130. B Pejova, M Najdoski, I Grozdanov, A Isahi. J. Mater. Sci.-Mater. Electron. 11:405,

2000.131. B Pejova, T Kocareva, M Najdoski, I Grozdanov. Appl. Surf. Sci. 165:271, 2000.132. M Izaki, T Omi. J. Electrochem. Soc. 144:L3, 1997.


133. D Raviendra, JK Sharma. J. Appl. Phys. 58:838, 1985.134. S Misra, HC Padhi. J. Appl. Phys. 75:4576, 1994.135. D Braunger, D Hariskos, T Walter, HW Schock. Sol. Energy Mater. Sol. Cells

40:97, 1996.136. D Hariskos, M Ruckh, U Ruhle, T Walter, HW Schock, J Hedstrom, L Stolt. Sol.

Energy Mater. Sol. Cells 41–2:345, 1996.137. R Bayón, C Mafftiotte, J Herrero. Thin Solid Films 353:100, 1999.138. J Herrero, MT Gutierrez, C Guillen, JM Dona, MA Martinez, AM Chaparro, R

Bayon. Thin Solid Films 361:28, 2000.139. LP Deshmukh, KM Garadkar, DS Sutrave. Mater. Chem. Phys. 55:30, 1998.140. M Skyllas-Kazacos, JF McCann, R Arruzza. Appl. Surf. Sci. 22/23:1091, 1985.141. NC Sharma, DK Pandya, HK Sehgal, KL Chopra. Thin Solid Films 42:383, 1977.142. NC Sharma, DK Pandya, HK Sehgal, KL Chopra. Mater. Res. Bull. 11:1109, 1976.143. NC Sharma, DK Pandya, HK Sehgal, KL Chopra. Thin Solid Films 62:97, 1979.144. NC Sharma, DK Pandya, HK Sehgal, KL Chopra. Thin Solid Films 157, 1979.145. GB Reddy, DK Pandya, KL Chopra. Sol. Energy Mater. 15:383, 1987.146. GA Kitaev, VF Markov, LN Maskaeva, LE Vasyunina, IV Shilova. Inorg. Mater.

26:202, 1990.147. GA Kitaev, LN Maskaeva, VF Markov, AY Kurkin, LE Vasyunina. Inorg. Mater.

25:1065, 1989.148. LN Maskaeva, VF Markov, GA Kitaev. Russ. J. Appl. Chem. 73:751, 2000.149. BB Nayak, HN Acharya, GB Mitra. Bull. Mater. Sci. 3:317, 1981.150. BB Nayak, HN Acharya. J. Mater. Sci. Lett. 4:651, 1985.151. YS Sarma, HN Acharya, NK Misra. Thin Solid Films 90:L43, 1982.152. RS Mane, CD Lokhande. Thin Solid Films 304:56, 1997.153. RC Kainthla, DK Pandya, KL Chopra. J. Electrochem. Soc. 129:99, 1982.154. GS Shahane, BM More, CB Rotti, LP Deshmukh. Mater. Chem. Phys. 47:263, 1997.155. GS Shahane, KM Garadkar, LP Deshmukh. Mater. Chem. Phys. 51:246, 1997.156. GS Shahane, LP Deshmukh. Mater. Chem. Phys. 70:112, 2001.157. I Grozdanov. Semicond. Sci. Tech. 9:1234, 1994.158. JM Doña, J Herrero. Thin Solid Films 268:5, 1995.159. SA Al Kuhaimi, Z Tulbah. J. Electrochem. Soc. 147:214, 2000.160. NC Sharma, RC Kainthla, DK Pandya, KL Chopra. Thin Solid Films 60:55, 1979.161. T Nakazawa, S Kuranouchi, A Ashida, N Yamamoto. In: 12th ECPV Solar Energy

Conf., Amsterdam the Netherlands, 1994, p 601.162. GK Padam, GL Malhotra, SUM Rao. J. Appl. Phys. 63:770, 1988.163. T Yamaguchi, Y Yamamoto, T Tanaka, Y Demizu, A Yoshida. Thin Solid Films

281–282:375, 1996.164. T Yamaguchi, Y Yamamoto, T Tanaka, A Yoshida. Thin Solid Films 344:516,

1999.165. D Hariskos, R Heberholz, M Ruckh, U Ruhle, R Schäffler, HW Schock. In: 13th

ECPV Solar Energy Conf., Nice, France, 1995, p 1995.166. KC Sharma, JC Garg. J. Phys. D: Appl. Phys. 23:1411, 1990.167. DS Sutrave, GS Shahane, VB Patil, LP Deshmukh. Mater. Chem. Phys. 65:298,

2000.


168. G Contreras-Puente, O Vigil, M Ortega-Lopez, A Morales-Acevedo, J Vidal, MLAlbor-Aguilera. Thin Solid Films 361:378, 2000.

169. LP Deshmukh, DS Sutrave, BM More, CB Rotti, KM Garadkar. Semicond. Devices421, 1996.

170. K Kushiya, T Nii, I Sugiyama, Y Sato, Y Inamori, H Takeshita. Jpn. J. Appl. Phys.35:4383, 1996.

171. PK Nair, L Huang, MTS Nair, H Hu, EA Meyers, RA Zingaro. J. Mater. Res.12:651, 1997.

172. B Mokili, Y Charreire, R Cortes, D Lincot. Thin Solid Films 288:21, 1996.173. GK Padam, SUM Rao. Sol. Energy Mater. 13:297, 1986.174. GK Padam, GL Malhotra, SUM Rao. Phys. Status Solidi (a) 109:K45, 1988.175. A Ennaoui, M Weber, M Saad, W Harneit, MC Lux-Steiner, F Karg. Thin Solid

Films 361:450, 2000.176. RN Bhattacharya. J. Electrochem. Soc. 130:2040, 1983.177. GK Padam. Mater. Res. Bull. 22:789, 1987.178. KR Murali. Thin Solid Films 167:L19, 1988.179. JC Garg, RP Sharma, KC Sharma. Thin Solid Films 164:269, 1988.180. PK Vidyadharan Pillai, KP Vijayakumar, PS Mukherjee. J. Mater. Sci. Lett.

13:1725, 1994.181. GN Chaudhari, S Manorama, VJ Rao. J. Phys. D: Appl. Phys. 25:862, 1992.182. GN Chaudhari, S Manorama, VJ Rao. Thin Solid Films 208:243, 1992.183. AR Patil, VN Patil, PN Bhosale, LP Deshmukh. Mater. Chem. Phys. 65:266, 2000.184. R Suarez, PK Nair. J. Solid State Chem. 123:296, 1996.185. VM Markov, LN Maskaeva, LD Loshkareva, SN Uimin, GA Kitaev. Inorg. Mater.

33:555, 1997.186. N Parhi, BB Nayak, BS Acharya. Thin Solid Films 254:47, 1995.187. M Ristov, G Sinadinovski, I Grozdanov, M Mitreski. Thin Solid Films 173:53,

1989.188. YF Nicolau. Appl. Surf. Sci. 22/23:1061, 1985.189. YF Nicolau, JC Menard. J. Cryst. Growth 92:128, 1988.190. YF Nicolau, M Dupuy, M Brunel. J. Electrochem. Soc. 137:2915, 1990.191. YF Nicolau, S Davied, F Genoud, M Nechtschein, JP Travers. Synth. Met. 42:1491,

1991.192. VP Tolstoy. Thin Solid Films 307:10, 1997.193. M Sasagawa, Y Nosaka. Phys. Chem. Chem. Phys. 3:3371, 2001.194. CD Lokhande, BR Sankapal, HM Pathan, M Muller, M Giersig, H Tributsch. Appl.

Surf. Sci. 181:277, 2001.195. S Lindroos, A Arnold, M Leskela. Appl. Surf. Sci. 158:75, 2000.196. MK Karanjai, D Dasgupta. Mater. Lett. 4:368, 1986.197. CD Lokhande, KM Gadave. Mater. Chem. Phys. 36:119, 1993.198. CD Lokhande, VV Bhad, SS Dhumure. J. Phys. D: Appl. Phys. 25:315, 1992.199. L Huang, PK Nair, MTS Nair, RA Zingaro, EA Meyers. J. Electrochem. Soc.

141:2536, 1994.200. Y Rodriguez-Lazcano, MTS Nair, PK Nair. J. Cryst. Growth 223:399, 2001.201. H Wilman. Proc. Phys. Soc. 60:117, 1948.202. SH Pawar, PN Bhosale. Mater. Chem. Phys. 11:461, 1984.


203. K Vaccaro, HM Dauplaise, A Davis, SM Spaziani, JP Lorenzo. Appl. Phys. Lett.67:527, 1995.

204. HM Dauplaise, K Vaccaro, A Davis, GO Ramseyer, JP Lorenzo. J. Appl. Phys.80:2873, 1996.

205. K Vaccaro, A Davis, HM Dauplaise, SM Spaziani, EA Martin, JP Lorenzo. J. Elec-tron. Mater. 25:603, 1996.

206. JH Schön, O Schenker, B Batlogg. Thin Solid Films 385:271, 2001.207. M Martin. In: J Cook ed. Passive Cooling. Cambridge, MA: MIT Press, 1989.208. Y Mastai, D Katzen, KD Dobson, G Hodes. Unpublished.209. GB Reddy, V Dutta, DK Pandya, KL Chopra. Sol. Energy Mater. 5:187, 1981.210. I Grozdanov, CK Barlingay, SK Dey. Mater. Lett. 23:181, 1995.


3Mechanisms of ChemicalDeposition

In this chapter, we critically discuss the various mechanisms for CD. Since the ki-netics of deposition are intimately related to the mechanism, this aspect is alsotreated here. This chapter is divided into several sections dealing with different as-pects of the mechanisms. However, there is a large degree of overlap of the mate-rial in these different sections; indeed, it was often not obvious in which section aspecific topic should be placed. While the main discussion of any specific pointwill be limited to the most relevant section, the same point will often be treated,more briefly, in other sections. However, for those readers who are looking for aspecific topic in this chapter, it would be prudent not to be confined to only whatappears to be the relevant section. For example, elements of the mechanism, prob-ably with different emphases, are dealt with also in the sections on nucleation andon kinetics, while many issues of nucleation are considered in the mechanisticstudies.

There is one example of a CD process (for deposition of tin sulphides) inwhich elemental sulphur dissolved in a nonaqueous solvent is used as a source forS. Since this appears to be the only example in the literature for this type of filmdeposition, it will be discussed in Chapter 6 together with the relevant study on tinsulphides. However, there is no reason to believe that this process may not be ap-plicable to other materials. Metal sulphides (and selenides) are known to form, asprecipitates, by reacting certain metal salts with dissolved elemental chalcogen, al-though visible film formation seems to be limited, up to now, to this one example.


3.1 THE FOUR BASIC MECHANISMS

In spite of the fact that CD has been in use for a long time and that the reactionsinvolved appear to be quite straightforward, the mechanism of the CD process isoften unclear. There is a good reason for this: There are several different mecha-nisms of CD. These can be divided into four fundamentally different types. Usingthe common thiourea deposition of CdS as an example, these are:

The simple ion-by-ion mechanism

Cd(NH3)42�D Cd2� � 4NH3

(dissociation of complex to free Cd2� ions)(3.1)

(NH2)2CS � 2OH� → S2� � CN2H2 � 2H2O

(formation of sulphide ion)(3.2)

Cd2� � S2� → CdS (CdS formation by ionic reaction) (3.3)

The simple cluster (hydroxide) mechanism

nCd2� � 2nOH�D [Cd(OH)2]n

(formation of a solid Cd(OH)2 cluster)(3.4)

[Cd(OH)2]n � nS2� [from Reaction (3.2)] →nCdS � 2nOH� (exchange reaction)

(3.5)

The complex decomposition ion-by-ion mechanism:

(NH2)2CS � Cd2�D [(NH2)2CSMCd]2� (3.6)

[(NH2)2CSMCd]2� � 2OH� → CdS � CN2H2 � 2H2O (3.7)*

The complex decomposition cluster mechanism:

[Cd(OH)2]n [from Reaction (3.4)] � (NH2)2CSD

[Cd(OH)2]n�1(OH)2CdMSMC(NH2)2(3.8)

[Cd(OH)2]n�1(OH)2CdMSMC(NH2)2 →[Cd(OH)2]n�1CdS � CN2H2 � 2H2O

(3.9)

which continues until conversion of all the Cd(OH)2 to CdS.

The first two mechanisms involve free sulphide ions (or other anions), whilethe last two are based on breaking of a carbon–chalcogen bond and do not involveformation of free chalcogenide. Most mechanistic studies have assumed the for-

* The use of the simple thiourea–Cd ligand in Eq. (3.7) is for simplicity. There are a number of dif-ferent complexes involving Cd and thiourea, in particular, some containing hydroxy groups. As longa a solid phase of Cd(OH)2 is not present, then such hydroxy–thiourea Cd complexes involve an ion-by-ion-type of mechanism, as exemplified in Eq. (3.7).


mation of free anions (not always justifiably). It is also possible that more than onemechanism occurs in parallel or that conditions change during deposition so thatthe deposition mechanism also changes.

Since most mechanistic studies were carried out for CdS deposition from anammoniacal solution and using thiourea as a sulphide source, and a smaller num-ber for CdSe deposition using selenosulphate as selenide source, this chapter willmostly use these materials as examples. However, it should be stressed that theconcepts discussed in this chapter are, for the most part, valid for the depositionof all semiconductors.

3.2 SLOW FORMATION OF THE CHALCOGENIDE IONS

The rate-limiting step in CD for the first two mechanisms is almost always for-mation of the chalcogenide ion. This reaction should be slow; otherwise fast, ho-mogeneous precipitation of the metal chalcogenide will occur with little film for-mation. (Even rapid precipitation can lead to a film; however, this film will beextremely thin and in most cases not visible.) Almost all the literature on CD islimited to sulfides (mostly), selenides, and oxides (including hydrated oxides andhydroxides). Anion-forming reactions are described in this section.

3.2.1 Sulphides

3.2.1.1 Thiourea

Thiourea (the sulphur analogue of urea) is the most commonly used sulphur pre-cursor. In alkaline solution (in which depositions involving thiourea are carriedout), the first step of the hydrolysis gives sulphide ions and cyanamide:

SC(NH2)2 � OH�D HS� � CN2H2 � H2O (3.10)

The cyanamide can hydrolyze further, with the overall reaction, if carried to com-pletion, going via urea to ammonium carbonate:

CN2H2 H2O→ CO(NH2)2 2H2O→ (NH4)2CO3 (3.11)

Cyanamide can also react with ammonia to give guanidine:

CN2H2 NH3→ (NH2)2CBNH (3.12)

All of these compounds can be (and have been) found as impurities in CD films.However, the important step is the formation of sulphide ion.

In neutral and acidic solutions, thiourea can be decomposed to thiocyanateion [1], which can be useful if the intention is to deposit thiocyanates:

SC(NH2)2 → NH4� � CNS� (3.13)


It should be kept in mind that many of these decomposition reactions areequilibria. The decomposition of thiourea in the absence of a metal ion will nor-mally be much slower than in the presence of such an ion. The metal ion removessulphide as metal sulphide—the less soluble the sulphide, the more effective theremoval at very low sulphide concentrations. This continuous removal of sulphideshifts the equilibrium to the direction of more sulphide production. The same prin-ciple holds for many other anion precursors.

3.2.1.2 Dimethylthiourea

Dimethylthiourea is much less commonly used than the more available thiourea:

(CH3)NHC(S)NH(CH3) � H2O → (CH3)NHC(O)NH(CH3) � H2S (3.14)

3.2.1.3 Thioacetamide

Thioacetamide has been used for a long time as an analytical reagent to precipi-tate metal sulphides (see Ref. 2 for many relevant references). Thioacetamide canbe hydrolyzed over a wide range of pH and is often used for CD in acidic baths.In a strongly acidic bath (pH � ca. 2), H2S is formed:

H3C.C(S)NH2 � 2H2O → CH3COOH � H2S � NH3 (3.15)

It has been shown [3] that this reaction can proceed by two pathways, one in whichthe carbon–sulphur bond is broken first:

H3C.C(S)NH2 � H2O → H3C.C(O)NH2 � H2S (3.16)

forming acetamide as an intermediate, or a pathway in which the carbon–nitrogenbond is first broken to give thioacetic acid:

H3C.C(S)NH2 � H2O → H3C.C(S)OH � NH3 (3.17)

which then is hydrolyzed to H2S and acetic acid. The H2S dissolves in water ashydrosulphide ion:

H2S � H2OD HS� � H3O� (3.18)

In an alkaline bath, the overall reaction is:

H3C.C(S)NH2 � 2OH� → CH3COO� � NH3 � HS� (3.19)

Hydrolysis in alkaline solution is considerably faster than in acid solutions.At intermediate pH values, particularly in weakly acidic solutions (pH � 2),

metal sulphide formation using thioacetamide may proceed through decomposi-tion of a metal ion (or solid phase)–thioacetamide complex rather than through in-termediate formation of sulphide [4] (see Sec. 3.3.3). Thioacetamide in pure wa-ter is fairly stable and does not readily hydrolyze at room temperature.


3.2.1.4 Thiosulphate

The earliest CD processes were carried out using thiosulphate. Although thiourea(and to some extent thioacetamide) are now more commonly used to deposit sul-phides, thiosulphate is still sometimes used. While the reaction pathways listedbelow are intended to suggest possibilities for reactions involving thiosulphate, itmust be noted that the mechanism(s) for these reactions is (are) still not clear.Mechanisms have been proposed in the CD literature, but no convincing proof forany particular one has been forwarded.

Thiosulphate depositions are most often carried out in a weakly acidic bath(pH 3). Several reactions are possible in such solutions:

S2O32� � H2O → H2S � SO4

2� (3.20)

with the H2S dissolving to give HS�, as in Reaction (3.18):

S2O32� � H� → S � HSO3

� (3.21)

S2O32� � 2H�→ S � SO2 � H2O (3.22)

and in alkaline solution:

S2O32� � OH� → HS� � SO4

2� (3.23)

It has been suggested that the thiosulphate, a reducing agent, may act as an elec-tron donor and reduce the elemental sulphur formed in Reactions (3.21) and(3.22), forming sulphide ions:

S � 2e� → S2� (3.24)

Because of the strong complexes thiosulphate forms with some metal ions,it is very possible that these metal–thiosulphate complexes undergo a complex-de-composition mechanism (Section 3.3.3). However, one early study on the forma-tion of PbS on boiling Pb2� and thiosulphate in water found that PbS formed morereadily when excess thiosulphate was present [5], which suggests that decompo-sition of thiosulphate to sulphide might be the dominant pathway under the con-ditions of that study.

3.2.2 Selenides

3.2.2.1 Selenourea

Selenourea (SeC(NH2)2) was apparently first used for CD by Kutzscher’s groupin Germany during World War II. It appears that this work was not published; ref-erences to it come through other sources [6]. The first published use of selenoureafor CD appears to be in 1949 by Milner and Watts [7] to deposit PbSe for photo-conductive cells. It was the main reactant used to form selenide films (mainlyPbSe) through the 1960s, after which selenosulphate (see later) became the dom-


inant selenide precursor. Other examples of its use are for ZnSe [8] and HgSe [9].Selenourea is an unstable compound that requires the presence of a reducingagent—usually Na2SO3—to minimize oxidation to elemental Se. The selenideformation presumably parallels thiourea hydrolysis in Eq. (3.10):

SeC(NH2)2 � OH� → HSe� � CN2H2 � H2O (3.25)

3.2.2.2 Dimethylselenourea

Dimethylselenourea (CH3)2NC(NH2)Se (for preparation see Ref. 10) was used todeposit films of PbSe [11]. It is more stable than selenourea, although still not verystable in solution and, like selenourea, was used together with Na2SO3 to preventoxidation. The optimum pH for this deposition was 9.8; bulk precipitation oc-curred at a pH of 10.1, while the deposition slowed greatly as the pH decreasedeven a few tenths of a pH unit. The activity of this solution, as measured by therate of deposition, was reported to increase with time (hours). This was explained,in view of the measured reduction in the rate of reaction by sulphite (which wasassumed to complex with the dimethylselenourea), by reduction of the concentra-tion of sulphite with time (by oxidation to sulphate). This means that, for repro-ducible results, the age of the dimethylselenourea/sulphite solution should betaken into account or an aged solution (at least 10 hr old) should be used (thechange in activity of the solution slows down to a large degree after this time).

3.2.2.3 Selenosemicarbazide

Selenosemicarbazide (H2NN(H)C(Se)NH2) was used by Velykanov et al. to makeCd, Zn, Ag, and Hg selenide films and precipitates from aqueous alkaline solu-tions [12]. As with the other Se precursors, Na2SO3 was used to stabilize the se-lenosemicarbazide against rapid decomposition. This reagent was reputed to bemore stable than the previous two.

3.2.2.4 Selenosulphate

Selenosulphate (Na2SeSO3), which is the analogue of thiosulphate with the activeS atom substituted by Se, was used by Kitaev for deposition of PbSe [13] andCdSe films [14]. Since it is more stable, simpler to prepare, and cheaper than se-lenourea, it simplified the deposition of selenides and has for a long time, withonly a few exceptions, been the precursor invariably used to deposit selenides. Itcan be prepared by dissolving elemental Se in an aqueous sulphite solution:

Se � SO32�D SeSO3

2� (3.26)

The Se dissolves slowly (typically an hour or two at 60–70°C). Although morestable than selenourea, it does slowly hydrolyze and becomes less active withage—fairly rapidly for the first few days after preparation and then slowly; it isstill usable for weeks after preparation without any special storage demands, al-


though at increasingly reduced activity. Storage under cold, oxygen-free condi-tions will slow the aging further. This gradual change in activity should be takeninto account in depositions using this material, since it is not always practical tomake it fresh for each deposition.

The mechanism for hydrolysis of selenosulphate is often given as

SeSO32� � OH� → HSe� � SO4

2� (3.27)

However, although SO42� is likely to be a final product, the reaction is probably

not as simple as this. A suggested first step for the hydrolysis is [15]

2SeSO32� � H2O → HSe� � SeS2O6

2� � OH� (3.28)

Selenosulphate is unstable in acidic solutions. If the pH of a fairly concen-trated selenosulphate solution is reduced below ca. 7, red Se will precipitate out.This property has been used to prepare films of Se by slightly acidifying dilute so-lutions of selenosulphate [16].

3.2.3 Tellurides

Very few examples of telluride deposition have been reported. There are a numberof reasons why tellurides are more difficult to deposit by CD than sulphides or se-lenides. One is the very negative redox potential of the telluride/tellurium couple(E0 � �1.14 V). This means that a strong reducing solution is necessary. Another,directly related reason is the instability of telluride ion (and possible tellurium pre-cursors). Even dissolved oxygen will rapidly oxidize telluride. The strong reduc-ing solution will need to reduce dissolved oxygen, preventing this reaction. Thus acombination of a suitable Te source and a strong reducing agent with exclusion (orreduction) of oxygen can succeed in forming tellurides. In fact, such conditionshave also been found necessary to form films of ZnSe using selenosulfate and arepreferable for deposition of ZnS films (see Sec. 3.5 and Chap. 4).

CdTe has been deposited by hydrazine reduction of TeO2 [17,18]. The po-tential of hydrazine oxidation is sufficiently negative to allow formation of tel-luride:

N2H4(aq) � 4OH� → N2 � 4H2O � 4e� E0 � �1.16 V (3.29)

Te � 2eD Te2� E0 � �1.14 V (3.30)

especially because only very small concentrations of telluride need to exist at anytime, causing a positive shift of the Te reduction through the Nernst equation (Eq. 1.32).

Te was considered to be essentially insoluble in Na2SO3 under normal con-ditions, although preparation of tellurosulphate, the analogue of selenosulphate, un-der hydrothermal conditions has been reported [19]. A recent study has describedCdTe deposition using this reagent [20]. Apparently the solubility, while low, is suf-


ficient to be useful. At this stage, it is assumed that the tellurosulphate reacts anal-ogously to selenosulphate (Eqs. 3.26–3.28). As with the TeO2 reduction just de-scribed, because of the great ease of oxidation of telluride to elemental Te, the filmsinvariably contain elemental Te. (See Sec. 4.3 for details of this deposition.)

While not true CD, a novel telluriding agent, Te dissolved in an alkaline so-lution of hydroxymethanesulphinic acid, has been used to convert Cd(OH)2 filmsto CdTe [21]. While there is some doubt as to the nature of the active telluridingagent, from the description of the preparation process of this reagent—a change incolor from deep purple (characteristic of polytelluride ions, Tex

2�) to a faint pink(pure Te2� is colorless but will be colored this way if traces of Tex

2� are present)as the preparation proceeds—it does appear to contain free telluride ion. It shouldbe noted that elemental Te can slowly dissolve in concentrated air-free alkalinesolutions, with the formation of polytelluride and its characteristic purple color.

3.2.4 Oxides

Oxides have been commonly deposited by CD. In many cases, the deposit is a hy-droxide or hydrated oxide formed by reaction of the metal ions with slowly gener-ated hydroxide. A variety of precursors has been used to generate the hydroxide.

3.2.4.1 Urea

Urea is slowly hydrolyzed in water to form ammonium carbonate:

(NH2)2CBO � 2H2O → (NH4)2CO3 (3.31)

Carbonates are alkaline since they dissociate to some extent to form hydroxideions:

CO32� � H2OD OH� � HCO3

� (K � 1.8 � 10�4) (3.32)

The hydroxide so formed will react with the metal ion to form the metal hydrox-ide, hydrated oxide, or oxide, depending on the relative stability of the various ox-ides and hydroxides. (The resulting hydroxides or oxy hydroxides can be heatedin air or oxygen to form the oxides.) In addition, insoluble carbonates may alsoform. The competition between hydroxide and carbonate will depend on their sol-ubility products. Carbonates tend to be more soluble than hydroxides of the samemetal ion. On the other hand, the value of K for equilibrium (3.32) (1.8 � 10�4)means that the concentration of hydroxide will be ca. four orders of magnitude lessthan that of carbonate (assuming no other pH-determining species is present).

While the urea method has been commonly used in the past to form bulkprecipitates of basic salts, including oxides, for analytical purposes, it was notedthat transparent, adherent films were typically formed on the walls of the beakerin which the precipitation was carried out [22]. Notably “basic stannic sulfate,”which was probably mainly SnO2 since very little sulphate was found in the ac-


companying precipitate, was found to adhere so tightly to the walls of the beakerthat it was difficult to remove [23].

3.2.4.2 Dimethylamineborane (DMAB) andTrimethylamineborane

Dimethylamineborane (DMAB) and trimethylamineborane have been used to de-posit Zn [24], In [25] and Fe [25a] oxides using the metal nitrates. The nitrate an-ion is believed to be reduced by the DMAB to nitrite and hydroxide:

(CH3)2NHBH3 � 2H2O → BO2� � (CH3)2NH � 7H� � 6e� (3.33)

NO3� � H2O � 2e� → NO2

� � 2OH� (3.34)

Mx� � xOH� → M(OH)x (3.35)

In the case of Zn the oxide forms spontaneously, while for In the hydroxide filmis heated to form the oxide. This mechanism is actually a mixed electrochemicalprocess involving partial anodic and cathodic reactions (i.e., charge exchange oc-curs at an interface) rather than a pure chemical deposition process.

3.2.4.3 Persulphate

Persulphate (also called peroxydisulphate) (S2O82�) is a very strong oxidizing

agent and has been used to deposit oxides, often in a higher oxidation state that thatof the original metal ion. Films grown using this reagent include �-PbO2, NiO,MnO2, and Tl2O3. The two latter required a small concentration of Ag� ions in thedeposition solution as a catalyst. Ag� is a known catalyst for oxidations usingS2O8

2� (it is oxidized by persulphate to Ag(III), which is then the active oxidationagent). It is probable that Pb and Ni also act likewise, while Mn and Tl do not (ormuch less so). A study of the use of persulphate for the deposition of PbO2 has pro-vided strong evidence that the deposition actually is a mixed electrolytic process(similar to that proposed for dimethylamineborane depositions) rather than a purechemical deposition [26]. The partial electrolytic reactions were given as

Pb2� � 2H2O → PbO2 � 4H� � 2e� (3.36a)

S2O82� � 2e� → 2SO4

2� (3.36b)

Persulphates hydrolyze to form (finally) sulphate and hydrogen peroxide; theprobable overall reaction can be given by

S2O82� � 2H2O → 2SO4

2� � H2O2 � 2H� (3.37)

This reaction occurs via various radical species, and it is also possible that someoxide depositions occur by these radicals (the formation of which is probably ini-tiated by fission of the S2O8

2� ion to two sulphate radicals) or directly by the hy-drogen peroxide, which itself usually involves a free-radical reaction. Hydrogenperoxide itself has been used directly to form oxide films in a few cases.


3.2.4.4 Other oxide-forming reactions

Aqueous solutions of acidic metal salts are usually inherently unstable and hy-drolyze readily to oxides (the hydroxides of these metals tend to be relatively un-stable), in some cases forming films. Such hydrolysis can be more readily con-trolled by adding boric acid to fluoro-complexes of the metal, e.g.:

H2TiF6 � 2H2OD TiO2 � 6HF (3.38a)

followed by removal of the HF by the boric acid:

H3BO3 � 4HFD BF4� � H3O� � 2H2O (3.38b)

Ref. 27, which describes such reactions as liquid phase deposition, gives more de-tails on this method.

Finally, many ammoniacal or amine solutions of metal salts will form filmsof hydroxides or oxides on heating or even long standing at room temperature.Ammonia is very volatile and will gradually be lost (in an open system), resultingin reduced complexation of the metal ions. The pH will also drop, but it is likelythat the increase in concentration of free metal ions due to loss of ammonia dom-inates. Chapter 7 treats the deposition of oxides in detail.

3.2.5 Halides

Halide ions can be formed by hydrolysis of alkyl halides, as shown for AgCl pre-cipitation [28]:

RMCl � OH� → Cl� � ROH (3.39)

Haloalcohols (halohydrins) which are more water soluble than alkyl halides, havebeen used to generate chloride ions [29], and are therefore more suitable for aque-ous CD. The hydrolysis of haloalcohols has been used to deposit films of AgBrand AgI, e.g.,

ICH2CH2OH � H2O → I� � HOCH2CH2OH � H� (3.40)

3.2.6 Other Anions

Although they do not appear to have been used (at least deliberately) to formfilms, there are other slow anion-generating reactions. Although most of the com-pounds formed in these reactions are not considered semiconductors in the nor-mally accepted sense, they merit at least a mention here.

3.2.6.1 Sulphate

Sulphate is formed by reaction between persulphate and thiosulphate:

S2O82� � 2S2O3

2� → 2SO42� � S4O6

2� (3.41)


This would be limited to cations that form insoluble sulphates but soluble persul-phates and thiosulphates (Ba and Sr were demonstrated by Lamer and Dinegar[30]). Thiosulphate, in particular, forms soluble complexes with many cations andtherefore should (often) not present a problem in this respect, as long as the metalsulphide is not formed under the conditions of the deposition. In addition, solventsother than water can be used in principle, and therefore it might be possible to de-posit sulphates that are soluble in water but insoluble in another solvent.

3.2.6.2 Phosphates

Phosphates have been formed by slow hydrolysis of trialkyl phosphates, hydrogenphosphate ions, or metaphosphoric acid, which liberate phosphate ions (see p. 46in Ref. 31). This may be of interest as a precursor for the preparation of phosphidesemiconductors.

3.2.6.3 Arsenates

Arsenates have been described in one case exploiting the fact that the zirconylcation forms a water-soluble arsenite but insoluble arsenate. By adding nitric acidto a solution of zirconyl chloride and sodium arsenite, the arsenite was oxidized toarsenate by the nitric acid, precipitating the insoluble zirconyl arsenate [32]. As forphosphates (and probably more readily), arsenates might be reduced to arsenides.

It should be stressed that these reactions were used to form precipitates andnot films. There is no guarantee that films can be formed using these reactions.However, it is reasonable to expect that, under the right conditions, it may be pos-sible to produce films of these compounds. It is left as an exercise for the curiousreader to find these “right” conditions.

Details of these and other slow precipitations are given in Ref. 31.There are a number of examples of homogeneous precipitation of hydrox-

ides based on slow cation release, such as destruction of the Fe–EDTA complexwith H2O2 (see Ref. 33). In CD, the only well-defined example of this is heatingan ammonia complex (e.g. of Cd2�). The loss of ammonia by volatilization willgradually increase the concentration of free Cd2� ions.

3.3 MECHANISMS OF CHEMICAL DEPOSITIONHaving dealt with slow formation of the reacting anion, we now turn to the variousmechanisms by which the CD compounds are formed. For the most part, the de-tails of nucleation and film growth are left to the following section. Here we con-centrate on the reactions taking place that form the semiconductor material. Thereare four main mechanisms for the compound formation, as outlined in Sec. 3.1;which one is operative depends on the specific process and reaction parameters.

1. Simple ion-by-ion mechanism2. Simple cluster (hydroxide) mechanism


3. Complex-decomposition ion-by-ion mechanism4. Complex-decomposition cluster mechanism.

3.3.1 (Simple) Ion-by-Ion Mechanism

The conceptually simplest mechanism, often assumed to be the operative one ingeneral, is commonly called the ion-by-ion mechanism, since it occurs by se-quential ionic reactions. The basis of this mechanism, illustrated for CdS, is givenby

Cd2� � S2�D CdS (adsorbed on the substrate) (3.42)

If the ion product, [Cd2�][S2�], exceeds the solubility product, Ksp, of CdS (ca.10�28; Table 1.1), then CdS can form as a solid phase, although a larger ionicproduct may be required if supersaturation occurs. If the ion product does not ex-ceed Ksp, no solid phase will form, except possibly transiently due to local fluctu-ations in the solution, and the small solid nucleii will redissolve before growing toa stable size. For that reason, the precipitation process is shown as an equilibriumrather than as a one-way reaction.

This reasoning, that no solid phase phase will form if the ion product doesnot exceed Ksp, may not be true under certain circumstances. CdS formation mightoccur at the substrate surface under conditions where none can be formed homo-geneously. This is an important point; however, since it will be treated in more de-tail in the following section, on hydroxide formation, it therefore will not be dis-cussed further in this section.

Equation (3.42) gives the overall reaction for formation of CdS. However,the process is more complicated than this and comprises a number of reactions andequilibria. The mechanism involves the formation of S2� ions and control of Cd2�

concentration.The S2� can be formed by a number of methods, as already described. Here

we consider the most common one, the decomposition of thiourea by aqueous al-kaline solution:

(NH2)2CS � 2OH� → S2� � CN2H2 � 2H2O (3.43)

Since the S2� concentration can be made as low as desired simply by controllingthe rate of Reaction (3.43) (e.g., using low temperatures and/or relatively low pH),in principle, even at relatively high free-Cd2� concentrations, the deposition rateshould be easily controlled. Since an alkaline pH is required to decompose thethiourea to sulphide, a complexing agent is needed to keep Cd2� in solution andto prevent Cd(OH)2 from precipitating. As explained in Chapter 1 [Eqs. (1.26) and(1.27) and calculation following these equations], using ammonia as a complex-ant for Cd (0.1 M total Cd concentration) and at a pH of 11, a concentration of am-monia � 1.19 M will be needed to prevent formation of Cd(OH)2 at room tem-


perature—more at the higher temperatures more commonly encountered in CdSdeposition. (We shall see, in the following section, that although this statement isperfectly accurate, it does not mean that CdS deposition will not occur at lowerconcentrations of ammonia at the same pH.)

When reading the literature, in many (probably most) cases it is not clearwhether the deposition proceeds by an ion-by-ion process. The reason is that, un-less another mechanism is specifically discussed, it is often assumed that the de-position proceeds via the ion-by-ion mechanism. If the exact deposition parame-ters are known, which mechanism is operative can, in most cases, be calculated.Two criteria have often been cited in the literature as proof of deposition via theion-by-ion mechanism. One is epitaxial deposition of the CD film. (Epitaxy refersto growth of one material on another in such a way as to result in coherence be-tween the lattice of the substrate and the deposit. Often—although not necessar-ily—the lattice of the deposit is aligned in the same direction as that of the sub-strate.) This is based on the expectation that a cluster mechanism will not result inan epitaxial film; for this to occur, clusters of maybe thousands of atoms wouldneed to be able to rearrange themselves on the substrate. Some examples of epi-taxial growth are given in Sections 3.4.2 and 4.1.5.2.

The second criterion, sometimes seen in the older literature, is that filmsformed via the cluster mechanism should be poorly adherent and optically scat-tering, while those formed via the ion-by-ion mechanism will be adherent andtransparent. Unlike the epitaxy, this criterion is faulty; films formed by the clus-ter mechanism can be highly transparent and strongly adherent, while there areexamples of films that proceed via the ion-by-ion mechanism that are not well ad-herent and are optically scattering. The degree of adhesion of the film and itstransparency say little, if anything, about the mechanism of the deposition. Thus,following this reasoning, it was claimed that ion-by-ion deposition occurred un-der conditions where a visible Cd(OH)2 suspension was present [34]. While thepresence of Cd(OH)2 does not exclude the possibility that ion-by-ion depositionoccurs, it is unlikely that ion-by-ion deposition would be dominant when a highconcentration of Cd(OH)2 is present.

More recently, there have been a small number of studies that provide strongevidence for the ion-by-ion mechanism. It must be pointed out, however, thatwhile it is not very difficult to distinguish between an ion-by-ion and clustermechanism in most cases, it is much more difficult to distinguish between a sim-ple ion-by-ion and a complex-decomposition ion-by-ion mechanism. Thereforemost investigations that conclude an ion-by-ion mechanism is operative, whileusually assuming the simple ion-by-ion process, do not distinguish between thesimple and complex-decomposition pathways.

One investigation has shown a clear-cut boundary in the crystal size of films(CdSe, CdS, and, to a lesser extent, PbSe), depending on whether the depositionoccurred via an ion-by-ion or a cluster mechanism [15]. The solution conditions


were constant except for a change in the ratio between complexant and Cd. Asmall crystal size was obtained under conditions where Cd(OH)2 was proven toexist (as a colloid, not visible to the eye) in the deposition solution, while a largercrystal size was obtained when no Cd(OH)2 was present as a solid phase. Thus theseparation between the conditions where the ion-by-ion and cluster mechanismsoccurred was clearly shown. This study is dealt with in more detail in the follow-ing section.

A recent investigation of CdS deposited from a thiourea bath found that theinitial deposit (on mica) formed as islands ca. 0.5 nm high and 10–40 nm across;further growth led to an increase in height without changing the lateral dimensions[35]. Such a growth mode supports an ion-by-ion mechanism, since a clustermechanism, whereby (presumably fairly symmetric) hydroxide clusters adsorb ona substrate, is expected to lead to an equally symmetric initial growth mode. Theisland size in this study was measured by atomic force microscopy, where the tipsize can determine the measured lateral dimensions of particles smaller than thetip. For this reason, scanning tunneling microscopy (or electron microscopy forlateral dimensions) would be more reliable in such measurements. This investiga-tion was carried out using triethanolamine as complexant and under conditionswhere deposition was slower than usual. As always the case, the results and inter-pretation cannot be extrapolated to other bath conditions.

3.3.2 Hydroxide (Cluster) Mechanism

We noted in the earlier example that Cd(OH)2 precipitation should be avoided andcalculated the concentration of ammonia required to prevent this precipitation. Inreality, CDs are quite often carried out under conditions where a metal hydroxide(or hydrated oxide) is formed. This might seem to imply that a precipitate of, e.g.,Cd(OH)2 is formed at the start of such depositions. In fact, this is (usually) not thecase; the Cd(OH)2 is formed, but either as a colloid rather than a precipitate, or asan adsorbed species on the substrate but not in the bulk of the solution. SinceCd(OH)2 is colorless and colloids usually do not scatter light (otherwise it istermed a suspension), this means that the Cd(OH)2 colloid is not apparent to theeye. The CdS is then formed by reaction of S2� ion with the Cd(OH)2:

Cd2� � 2OH� → Cd(OH)2 (3.44)

followed by

Cd(OH)2 � S2� → CdS � 2OH� (3.45)

The driving force for Eq. (3.45) is the much lower value of Ksp for CdS (ca. 10�28;Table 1.1) than for Cd(OH)2 (2 � 10�14), which reflects the more negative freeenergy of formation of the former. This means that sulphide will readily substitutefor hydroxide in the case of Cd. An idea of the amount of sulphide needed to con-


vert Cd(OH)2 into CdS can be calculated from the ratio of the two solubility prod-ucts:

�

� � � 5 � 10�15(3.46)

At a pH of 11 ([OH�] � 10�3 M at room temperature), this gives a value for [S2�]of 5 � 10�21 M. Thus it requires a very low concentration of S2� indeed to beginto convert Cd(OH)2 into CdS. Even taking into account that most of the sulphidespecies will be HS� rather than S2�, at pH 11 the [S2�]:[HS�] ratio is 10�6.3 (or5 � 10�7) [Eq. (1.17)], and the required HS� concentration will be

� 10�14 M (3.46a)

still a very low concentration. At the higher temperatures usually used for CdS de-position, the OH� concentration will be about an order of magnitude higher (be-cause of the strong temperature dependence of the ionic product of water), and therequired S2� concentration will be about two orders of magnitude higher than re-quired at room temperature.

The participation of Cd(OH)2 in the deposition of CdS (and other metalchalcogenides) has been demonstrated or suggested on many occasions. Kitaev etal. presented a theoretical thermodynamic treatment of the Cd2�/ammonia/thiourea system to show when Cd(OH)2 should be present as a solid phase in thedeposition solution [36]. A graphical representation of this analysis is shown inFigure 3.1. This graph is based on two equilibria: the solubility product ofCd(OH)2 and the stability constant of the ammonia (ammine) complex of Cd.Consider first the former:

Cd2� � 2OH�D Cd(OH)2 Ksp � 2 � 10�14 (3.47)

Based on this equilibrium, we can express [Cd2�] in terms of [OH�] for the casewhere Cd(OH)2 will just precipitate; at higher pH it will precipitate and at lowerpH it will not. Since [OH�] can be converted into [H�] (and therefore pH) throughthe ion product of water (see Chap. 1), a graph can be made of pH vs. [Cd2�] (orp[Cd2�], which, analogously to pH, is equal to minus the logarithm of [Cd2�]).This is the hydroxide line in Figure 3.1. Its physical meaning is that above thisline, Cd(OH)2 will be present in the solution, while below it there will be noCd(OH)2.

A similar (somewhat more complicated) calculation can be made based onthe stability constant of the Cd–tetrammine complex [see Eq. (1.27)] and using thehydrolysis of ammonia (see Refs. 34 and 36 for details of the calculation), which

5 � 10�21

��5 � 10�7

10�28

��2 � 10�14

[S2�]�[OH�]2

[Cd2�][S2�]��[Cd2�][OH�]2

Ksp(CdS)��Ksp(Cd(OH)2)


gives the free-Cd2� concentration as a function of pH. This is shown as the com-plex line in Figure 3.1. (The calculation is made for a total-Cd2� concentration of0.1 M, but it is only slightly dependent on this concentration; e.g., for a Cd2�

concentration of 1 mM, the line will shift ca. 0.2 pH units to lower pH values.Also, the presence of other ammine and hydroxy complexes should strictly speak-ing be taken into account; however, these considerations will not result in majorchanges to the overall picture.) Since the free-Cd2� concentration is a functionof ammonia concentration, the p[Cd2�] can be also identified with a p[NH3](top axis).

Considering both the hydroxide and complex lines together, for the left sideof the figure, where the complex line is above the hydroxide line, the concentra-tion of free Cd2� will always be high enough to form Cd(OH)2. Where the com-plex line is below the hydroxide line, however (the right side of the figure),Cd(OH)2 will not form at pH values above the complex line but below the hy-droxide line, since the Cd2� concentration will not be high enough to form

FIG. 3.1 Regions of stability for the Cd–ammonia system for 0.1 M total Cd concentra-tion and at room temperature. The hydroxide line separates conditions where Cd(OH)2 will(above the line) or will not (below the line) thermodynamically exist. The complex linegives the concentration of free Cd2� at any pH value (where the pH is determined only bythe ammonia concentration and not, e.g., by added alkali). (Adapted from Ref. 34.)


Cd(OH)2. Extra OH� ions (e.g., added KOH) are required to increase the pH andform Cd(OH)2 in this region.

This type of analysis (which can be extended to other metal ions and othercomplexants) provides a basis for choosing conditions for depositions involvinghydroxide clusters. An elegant example of this was given by Kitaev et al [36] forCdS deposition in the absence of ammonia (or any other complex, except forthiourea itself, which is a weak complexant for Cd). Extrapolation of the hydrox-ide line to p[Cd2�] � 1 (0.1 M free Cd2�) gives a pH of 7.7, or, alternatively, apOH of 6.3; above this value of pOH, no Cd(OH)2 is formed. However, at highertemperature, the ion product of water increases considerably (Table 1.2). Thus,while pOH (� pH for pure water) at room temperature is 7, at 80°C it becomes6.3. Thus, while no Cd(OH)2 is formed in a 0.1 M solution of a Cd salt (ignoringeffects of the Cd salt and thiourea on the solution pH), at 80°C it can thermody-namically form. Kitaev et al. demonstrated experimentally that a solution of 0.1M cadmium acetate and thiourea produced a film of CdS only at 90°C and above.This analysis, of course, is based on the hydroxide cluster model. Deposition byanother mechanism could conceivably occur at lower temperatures given enoughtime.

Rieke and Bentjen made a detailed study of the role of Cd(OH)2 in the de-position of CdS films from the ammoniacal thiourea bath [37]. They first studiedCd(OH)2 formation in the absence of thiourea. Cd(OH)2 formation, as measuredvisually by laser scattering (which shows up small amounts of turbidity in the so-lution), began at a pH of ca. 10.4 (the solid phase, once formed, could exist forquite long periods of time down to a pH as low as 10). Surface analysis (XPS)showed that Cd(OH)2 formed on silicon substrates at a pH value as low as 9, eventhough no Cd(OH)2 formed in the solution under those conditions. The Cd(OH)2,about 1.5 nm average thickness, was apparently stabilized against dissolution bythe substrate. In the presence of thiourea, a SEM study of the early stages of de-position showed that large, platelike crystals of CdS formed at a pH whereCd(OH)2 was present in the solution but that this deposit was very nonadherentand could easily be wiped off. However, at lower values of pH, where Cd(OH)2

had previously been shown to occur only on the substrate, strongly adherent CdSspherules were observed. The density (in terms of surface coverage) of this de-posit decreased with decreasing pH, and the coverage became low at low valuesof pH, where surface-adsorbed Cd(OH)2 did not occur. Therefore, only films de-posited under conditions where appreciable amounts of surface-adsorbed, but notbulk, Cd(OH)2 occurred were of good quality (adherent and specularly reflecting).

In connection with this study, in particular the suggestion that the Cd(OH)2

was stabilized by the substrate against dissolution, it has been shown thatCo(OH)2 can form at a solid (SiO2) surface at a pH lower than that necessary tocause bulk precipitation of Co(OH)2 [38]. This was explained by the effect of theelectric field at the solid/liquid interface on the dielectric constant of the interface


region and, through this, on the free energy (and therefore the solubility product)of the relevant hydroxide dissolution. It was noted that this field effect should in-crease with an increase in the charge of the cation. In terms of CD, this wouldmean that the difference in the ease of deposition of the hydroxide at a solid sur-face and homogeneously in solution will increase with the charge on the cation.Furthermore, this argument is not limited to hydroxides but should be valid for alldepositions, including chalcogenides (and therefore also ion-by-ion deposition),as was briefly mentioned in the previous section.

Betenekov et al. [39] used an isotopic tracer technique to show that, for theirrange of solution compositions, the initial deposition involved adsorption ofCd(OH)2 on the glass substrate. At the beginning of the reaction, only Cd was ob-served to form on the substrate and this was interpreted to be due to Cd(OH)2,since any other insoluble Cd compounds that might be formed from the depositionsolution (containing CdCl2, NaOH, NH4OH, and thiourea dissolved in water)were expected to contain either S or C.* However, they concluded that the depo-sition proceeded, not by reaction between Cd(OH)2 and sulphide formed by de-composition of thiourea, but rather by decomposition of a Cd(OH)2–thioureacomplex (see Sec. 3.3.3.1).

O’Brien and Saeed studied the CdS deposition, only using ethylenediamineinstead of ammonia as complexant because of the better-defined coordinationproperties of the former [40]. In common with the former studies, they also foundthe presence of Cd(OH)2 necessary for the deposition of good-quality CdS filmsby comparing the conditions needed to obtain such films with those calculated us-ing the relevant thermodynamic parameters.

A short digression at this point is required to define the term good-qualityfilm. The use of the adjective good-quality depends very much on what is requiredfrom the film. In the context of CD and thin-film preparation in general, goodquality almost always refers to two parameters: good adhesion and specular re-flection (smooth). Of course, requirements can be envisaged where these proper-ties, in particular the latter, are not desired, such as if a rough “scattering” film isrequired. However, keeping this caveat in mind, we will continue to use the termgood-quality as it is almost always used in the CD literature.

While most mechanistic investigations have been carried out on CdS, othersemiconductors, in particular CdSe, have also been studied with regard to the de-position mechanism. Kainthla et al., in their study of the formation of CdSe filmsfrom ammoniacal solutions containing sodium selenosulfate, noted that when avisible precipitate of Cd(OH)2 was present in their solutions (obtained by adding

* It needs to be repeatedly emphasized that the mechanism deduced in any one investigation is not nec-essarily valid for different experimental conditions. At the same time, it does appear likely that for thisvery well-studied system of CdS deposition from ammonia/thiourea baths, Cd(OH)2 is, in many cases,involved in the deposition.


KOH), a thin terminal thickness (ca. 80 nm) of CdSe was obtained, compared withseveral hundred nanometers obtained under optimal conditions (no precipitate)[41]. If the pH was too high and Cd(OH)2 precipitation was extensive, no film wasobserved to form. In fact, an analysis of their results suggests that films are ob-tained only if Cd(OH)2 is present (not necessarily as a visible precipitate). Whilethey do not specifically consider Cd(OH)2 as a chemical intermediate in their re-action, they do conclude that Cd(OH)2—both adsorbed on the substrate and in thesolution—acts as a nucleation center for CdSe formation.

Gorer and Hodes carried out a study of CdSe deposition from selenosul-phate solutions of Cd complexed with nitrilotriacetate—NTA [15], which is astronger complex for Cd than is ammonia and with which is easy to obtain condi-tions where no Cd(OH)2 is present, even at relatively high values of pH. A changein reaction mechanism, from hydroxide cluster to ion-by-ion, was observed bymonitoring the optical transmission spectra of the films deposited on glass. Thebasis for this investigation is the change in semiconductor bandgap, and thus in itsspectrum, when the crystal size becomes very small—the quantum size effect(discussed in detail in Chap. 10). The CdSe crystal size (and therefore thebandgap) is not strongly dependent on deposition parameters within a fairly widerange of parameters. However, it was observed that if the ratio between the NTAand Cd concentrations was increased above a certain value—called the criticalcomplex ratio—the crystal size suddenly increased (for fixed values of tempera-ture and pH). This was seen first by a red shift in the optical spectrum (see Chap.10 for examples of these changes) and was subsequently verified by direct mea-surements of the CdSe crystal sizes in the films using TEM and XRD.

A variety of analytical techniques were then used to verify that Cd(OH)2

was present in the solution when the complex:Cd ratio was below the critical value(Rc) but absent above it. Cd(OH)2 absorbs in the UV range of the spectrum, andspectral monitoring of Se-free solutions showed that it was present only below Rc.Light scattering by a blue laser also confirmed the presence of a heterogeneousphase below Rc but not above it. Similar XPS analyses to those employed by Riekeand Bentjen for CdS showed that Cd adsorbed on the glass substrate, immersed inSe-free solutions, only below Rc. This is seen in Table 3.1: Appreciable amountsof Cd (as Cd(OH)2) were seen only when the pH was sufficiently high and thecomplex:Cd ratio relatively low.

TABLE 3.1 Cd:Si Ratio Measured by XPS

pH 9.0 10.0 10.5 11.0 11.0Cd:Si 0.020 0.024 0.17 0.45 0.005

For glass substrates immersed in Cd2�/nitrilotriacetate solutions (no selenosulphate) at different pHvalues for 5 min at 40�C. NTA:Cd ratio � 1.63 except for the rightmost column, where it is 2.25.


Theoretical thermodynamic calculations of the conditions under whichCd(OH)2 should form were also carried out, based on the solubility product ofCd(OH)2, the stability constants of the Cd-NTA system, and the ion product ofwater at different temperatures. The values of Rc computed from these calcula-tions agreed with those measured experimentally.

In agreement with Kainthla et al. and contrary to at least some conditions forthiourea-based CdS deposition, Gorer and Hodes found that adsorption of a Cdhydroxide species on the substrate occurred only under conditions where the solidhydroxide formed also in the solution. This need not necessarily be interpreted ascontradictory; it may be due to the different conditions involved.

An important factor in the reaction for CdSe deposition via the hydroxidemechanism was an observed gradual increase in pH, ca. 0.8 pH units over thecourse of the deposition. It is not clear what the cause of this increase is. However,it means that while Cd(OH)2 may not be formed at the start of the deposition, itmay form during the deposition. This can then explain the induction time whereno apparent reaction takes place initially. For the preceding experiments, the stan-dard conditions were to set the solution pH to 10. Coloration of the solution (in-dicating formation of CdSe) occurred when the pH reached ca. 10.3 (depending,of course, on the temperature and complex:Cd ratio). If the pH was adjusted to10.3 at the beginning of the reaction, coloration began almost immediately ratherthan after a more typical time of several minutes. This increase in pH did not oc-cur for the ion-by-ion mechanism, probably because the excess NTA necessary inthis case also acts as a buffer. In any case, an increase in pH is not required for thismechanism.

These results for CdSe were extended to the deposition of PbSe (using cit-rate as complex) and CdS (using NTA and thiourea) and found to be also applica-ble for these cases [15]. For PbSe, the transition, while clearly evident, was not sosharp, since citrate is a weak complex and a solid phase is always present, but togreatly differing degrees depending on the conditions. The colloidal phase in thiscase is a hydrated lead oxide, which, in a selenosulphate-free solution, adsorbsstrongly onto the substrate (see Sec. 3.4.3 and Fig. 3.5). This hydrated oxide israpidly selenized to PbSe, and the process involves breaking down of the rela-tively large oxide crystals during reaction with the selenosulphate and recrystal-lization of the PbSe product. This can be seen by following the reaction of pre-cipitated hydrated lead oxide with selenosulphate by XRD (Fig. 3.2), where thesharp peaks of the oxide transform into very broad peaks indicative of an almostamorphous structure, which themselves vanish as sharp PbSe peaks appear.

A quartz crystal microbalance study of the kinetics of CdSe deposition fromthe preceding solution (nitrilotriacetate) showed an importance difference in themode of growth of the CdSe below and above Rc [42]. Below Rc (hydroxide mech-anism), a periodicity in the deposition rate was observed, with a period corre-sponding to a film thickness of ca. 5 nm—i.e., the approximate size of a single


crystal. Above Rc, the period corresponded to 0.3 nm, a single CdSe monolayer,as expected from an ion-by-ion growth. This interpretation is based on growth bycoverage of the surface by one layer—whether a layer of crystallites or a singleCd-Se layer—at a time.

The critical ratio concept already described was derived for a fixed concen-tration of metal ions. As pointed out by O’Brien and McAleese [43], while usefulat high metal concentrations, it requires modification when low concentrations areemployed, due to the changes in stability of complexes when dilute. They devel-oped a system whereby, on a plots of total metal concentration against total com-plex concentration, a curve defining a constant value of free-metal concentrationwas drawn. Figure 3.3 shows several such plots for the Cd–ethylenediamine sys-tem studied by them. From these plots, the need for much higher complex:metalratios for dilute solutions is evident. The complex:metal ratio required for a con-stant concentration of free metal, shown in the top curve, is reasonably constantfor higher concentrations but increases strongly at low concentrations.

We postpone detailed explanation of what determines the different crystalsizes in the two mechanisms to Section 3.4. At this point, it is enough to say thatthe CdSe crystal size in the hydroxide mechanism will be determined mainly bythe size of the Cd(OH)2 particles in the solution and on the substrate, while that

FIG. 3.2 XRD spectra showing the process of PbSe formation from the reaction of pre-cipitated hydrated lead oxide with Na2SeSO3 solution. (a) Starting material; (b–e) after 1.5,3, 4.5, and 6 mn reaction, respectively. (Adapted from Ref. 46.)


formed in the ion-by-ion mechanism will depend on the heterogeneous nucleationon the substrate.

It should be repeated at this stage that this mechanism is dependent on alarge difference between the solubility products of the hydroxide and chalco-genide of the required metal. The situation for ZnS, for example, is considerablyless favorable. The values of Ksp for the hydroxide and sulphide of Zn are 8 �10�17 and 3 � 10�25, respectively. The same calculation for Zn as carried out ear-lier for CdS shows that a S2� concentration of 4 � 10�15 M is required at pH �11 to convert the hydroxide into sulphide; this is a million times more than that re-quired to form CdS. This does not tell us at this stage whether ZnS will form ornot, but only that it is less likely to than CdS. We will return to this problem whenwe discuss the specific deposition of II–VI compounds in Chapter 4.

It is quite possible that the mechanism will change in the course of the depo-sition. As the metal is depleted from solution, the complex:metal ratio will increase

FIG. 3.3 Equivalent solution contour plots for solutions of Cd2� and ethylenediamine[en] at 50°C. The curves represent conditions for a constant concentration of free Cd2� �10�9M. Bottom curve (solid line): [total Cd] against [ethylenediamine]. Top curve (dottedline): the ratio of [ethylenediamine] to [total Cd] against [ethylenediamine]. (Adapted fromRef. 43.)


and may pass the point where no solid hydroxide phase is present in the solution.In this case, the ion-by-ion process will occur if the conditions are suitable. This ismost likely if the initial conditions were close to the border between ion-by-ion andcluster mechanisms. The opposite may also occur: initial deposition by an ion-by-ion pathway followed by clusters, which either build up gradually in the solutionor change their aggregation properties, adhering to the film. This has been shownto occur for CdS deposited from a thiourea bath using a combination of quartz crys-tal microbalance (which measures the mass of the deposit) and electrochemical im-pedence spectroscopy (which provides indirect structural information) [44].Change in aggregation properties of the colloids present during the deposition wassuggested as the cause of the change in deposition. It is also possible that bothmechanisms occur in parallel. Thus, deposition might begin with nucleation ofCd(OH)2, but growth might occur on this Cd(OH)2 via an ion-by-ion mechanism:

Cd(OH)2 � S2� � Cd2� → Cd(OH)2�CdS (3.48)

either by itself or together with conversion of the Cd(OH)2 to CdS, as in Eq.(3.45). Alternatively, each mechanism may occur independent of each other. Anexample of where this latter case appears to occur is in the deposition of PbSe[45,46], where, under certain conditions, two spatially separated domains of smallcrystals (cluster mechanism) and larger crystals (typical of the ion-by-ion mecha-nism) are obtained. The relative concentrations of these two domains can, ofcourse, be varied by changing the solution composition, but the fact that they oc-cur as separated domains on the substrate suggests that differences in the substratefrom one region to the other also plays a role. Another factor that could lead to achange in the details of the mechanism with time is the buildup of reaction prod-ucts in the solution. For example, the cyanamide formed in the decomposition ofthiourea [Eq. (3.10)] can react with Cd ions adsorbed on the surface of the CdS togive the sparingly soluble CdCN2:

(CdS)n Cd2�ads � CN2H2D CdS�CdCN2(ads) � 2H� (3.49)

The CdCN2 would then react with any sulphide present to form the more insolu-ble CdS. This would be parallel to the more straightforward sulphidization of(CdS)n Cd2�

ads to CdS.Before finishing this section, we present a few words concerning the effect

of deposition temperature on the terminal film thickness. Although exceptionshave been reported, it is a general observation that a lower deposition temperatureresults in ultimately thicker films (of course, the deposition rate is slower). Theremay be a number of reasons for this. One is that higher temperatures favor the hy-droxide cluster mechanism due mainly to the higher hydroxide concentration (thestrongly temperature-dependent ion product of water) and also to the lower sta-bility of complexes at high temperature (which may be offset by higher solubilityof the metal hydroxide). This effect has been shown for CdSe deposition, where a


higher complex:Cd ratio is needed to prevent Cd(OH)2 formation at higher tem-peratures [15]. The ion-by-ion mechanism, since it often does not result in homo-geneous precipitation, or at least less than in the cluster mechanism, tends to yieldfilms of larger terminal thickness. This will be valid even if the mechanism is amixture of the two mechanisms.

Another effect of low temperature is (in the absence of stirring) reducedmass transport. This can cause two opposing effects: reduction of reactants reach-ing the substrate and reduction of aggregation (collisions between colloidal parti-cles). The concentration of free chalcogenide ions will also increase essentiallyexponentially with increasing temperature. However, it is difficult to predict theeffect this will have on terminal thickness; rates of homogeneous precipitation andfilm formation will both increase. It is possible that thinner films will occur for theion-by-ion mechanism, since there is an increased probability of the occurrence ofhomogeneous precipitation with increasing chalcogenide ion concentration (theproduct of chalcogenide concentration and metal ion will increase—also becausethe free-metal ion concentration will be somewhat higher at higher temperaturedue to reduced stability of the metal complex).

3.3.3 Complex-Decomposition Mechanism

The complex-decomposition mechanism can, as with the free-anion mechanisms,be divided into ion-by-ion and cluster pathways. However, since experimentaldata relating directly to the complex decomposition mechanism is rather sparse, itwill be dealt with in one main section rather than two. The cluster pathway hasbeen more emphasized in these studies and will be dealt with first. It is very im-portant to stress at this point that almost all experiment data described below couldbe explained in terms of a simple (anion-mediated) mechanism. Equally valid,most of the data described in the previous section could be explained by a com-plex-decomposition mechanism.

3.3.3.1 Cluster Mechanism

The basis of this mechanism is that a solid phase is formed but, instead of react-ing directly with a free anion, it forms an intermediate complex with the “anion-forming” reagent. Continuing with CdS deposited from a thiourea bath as our ex-ample, this would be given as

MCd(OH)2 � (NH2)2CSD Cd(OH)2�SC(NH2)2 (3.50)

whereMCd(OH)2 is one molecule in the solid-phase cluster. This complex, or asimilar one containing also ammine ligands, then decomposes to CdS:

MCd(OH)2�SC(NH2)2 →MCdS � CN2H2 � 2H2O (3.51)

i.e., the SMC bond of the thiourea breaks, leaving the S bound to Cd.


Such a mechanism was suggested by Betenekov et al. based on their isotopictracer technique discussed in the previous section [39]. They suggested thatCd(OH)2 forms initially on the substrate and catalyzes the thioureau decomposi-tion. Of course the catalytic effect of the solid surface could be to decomposethiourea to sulphide ion and not necessarily to catalyze the complex-decomposi-tion mechanism.

A similar catalytic effect of PbS on the decomposition of thiourea had beensuggested previously by Norr [47]. Kinetic measurements by Rieke and Bentjensuggested that CdS likewise catalyzed thiourea decomposition [37]. Ortega-Borges and Lincot also deduced such a mechanism based on kinetic measurementsof the CdS deposition using a quartz crystal microbalance [48]. In this case, themeasurements were found to fit best with a complex-decomposition model. Boththey and Rieke and Bentjen found optimum deposition to occur under conditionswhere Cd(OH)2 was formed as a surface species on the substrate but not in the bulkof the solution. Kinetic measurements also led Doña and Herrero to a similar con-clusion of a complex-decomposition mechanism, but with the main difference thatthe initial adsorbed species is not Cd(OH)2 itself but an ammine–hydroxide [49]:

Cd(NH3)42� � 2OH�D [Cd(OH)2(NH3)2]ads � 2NH3 (3.52)

They based this modification on the known adsorbance of OH� on glass and onthe common occurrence of transition metal mixed water–ammonia complexeswith coordination number of 4. Parallel structural studies of the deposited CdSshowed textured growth, supporting a mechanism whereby alternate Cd and Sspecies were involved, in an ion-by-ion process. Such a growth suggests adsorp-tion of a molecular hydroxy-ammine species rather than a cluster. In fact, themechanism of Ortega-Borges and Lincot also does not differentiate between a hy-droxide cluster and molecule.

Unfortunately, it is nontrivial to distinguish reliably between the complex-decomposition and sulphide-formation mechanisms. For example, in the studyof PbS (as a precipitate) formation from thiourea [47], the two main results usedto support complex decomposition were: (a) very little sulphide was formed inalkaline solutions of thiourea and (b) addition of PbS powder catalyzed the re-action, seen by the disappearance of the induction time for precipitation andmore rapid PbS formation when PbS was added at the start of the reaction. How-ever, these results would also be obtained in a free-anion mechanism, for the fol-lowing reasons:

(a) Thiourea decomposition is an equilibrium reaction [see Eq. (3.10)].Formation of sulphide will shift the equilibrium back to the left. If ametal ion, which forms a sulphide with a low-solubility product, is pre-sent in the solution, however, it will remove even a very low concen-tration of sulphide continually as formed.


(b) The presence of a large solid surface in the solution will reduce the in-duction time, even if the mechanism proceeds through free-sulphideformation, since the initial nucleation step will be facilitated.

These factors do not argue against the complex-decomposition mechanism, butthey should not be too readily interpreted, in the absence of other evidence, as ev-idence against the sulphide mechanism. Granted, this is an old study, but it doespoint up the difficulty in distinguishing between the two mechanisms. Kineticstudies and subsequent fitting of the data from these studies to various models[48,49] appear to be the best way of approaching this problem at present.

3.3.3.2 Ion-by-Ion Mechanism

Consider the complexation of free Cd2� by thiourea to give a Cd–thiourea com-plex ion:

Cd2� � (NH2)2CSD [(NH2)2CSMCd]2� (3.53)

This ion could, in principle, hydrolyze by breaking the SMC bond to form CdS:

[(NH2)2CSMCd]2� � 2OH� → CdS � CN2H2 � 2H2O (3.54)

This would lead to CdS formation in solution.If the Cd2� is adsorbed on the substrate (either directly or indirectly through

a hydroxide linkage) or on previously deposited CdS, then the same reactionwould occur. If the CdS so formed remained bound to the substrate (it is assumedthat CdS generated on previously deposited CdS would remain bound), the resultwould be film growth by an ion-by-ion, complex-decomposition mechanism. Aswith the cluster mechanism, it is difficult to distinguish experimentally betweenthe complex-decomposition mechanism and the free-anion pathway.

Some studies have involved deposition (or precipitation) from acidic solu-tions. It is reasonable to assume that no hydroxide is present under these condi-tions for most metal ions commonly used in CD and that deposition occurs via anion-by-ion mechanism.

Thioacetamide decomposition at intermediate pH values, particularly inweakly acidic solutions (pH � 2), has been suggested to occur through a thioac-etamide complex rather than through intermediate formation of sulphide [4]. Ofcourse, this process may also occur in parallel with either acid or alkaline hydrol-ysis of thioacetamide to (ultimately) sulphide at certain pH ranges. It is also pos-sible that this complex-decomposition reaction occurs at both high and low pHvalues in certain cases.

The Cd–thiourea example that has been mainly used up to now in this sec-tion is a very weak complex. However, there are examples where the chalcogenideprecursor is a strong complexant to the metal and may also be used as the com-plexant. Depositions based on thiosulphate as a S source are good examples of


this. Also, thiosulphate depositions are most often carried out under at least some-what acidic conditions, which means that they are more likely to involve an ion-by-ion mechanism. Two old examples are the hydrolysis of silver thiosulphate togive (bulk) silver sulphide [50]:

Ag2S2O3 � H2O → Ag2S � H2SO4 (3.55)

and the deposition of a number of metal sulphides, including PbS and CuxS byheating thiosulphate solutions of the metal salts [51]. While less common nowa-days as a S source than, e.g., thiourea, thiosulphate reactions have been used to de-posit films of many different metal sulphides. The mechanism often suggested inthese studies is reduction of elemental S, formed by acidic decomposition of thio-sulphate, by the thiosulphate itself, forming sulphide ions (see Sec. 3.2.1.4). How-ever, no mechanistic studies of these reactions appear to have been undertaken.While there is no convincing proof in the literature to distinguish which mecha-nism is operative in these cases (sulphide-ion formation or complex decomposi-tion), chemical intuition leads us to expect the latter where the metal–chalcogenbond is strong (metal ions as Ag� and Cu� form very strong complexes throughthe labile S atom of thiosulphate), and it seems reasonable to expect thismetal–sulphur bond to remain intact during the reaction rather than formation ofsulphide ion to occur. At the same time, those same metals from sulphides withvery-low-solubility products (understandably, since both the solubility productand strength of complexation are related in the same way to the strength of themetal–sulphur bond). Therefore, very low concentrations of free sulphide areneeded to form the metal sulphide.

While thiosulphate is not very commonly used to deposit sulphides nowa-days, its Se counterpart, selenosulphate, is the most common reagent for selenidedeposition. By analogy with thiosulphate, it might be argued that the mechanisminvolves formation of selenide either through hydrolysis or through reduction ofSe, which forms even more readily in selenosulphate than does S in thiosulphate,or by complexation with metal ion or metal hydroxide and breaking of the SeMSbond (complex-decomposition mechanism). It can only be said that the mecha-nism of selenide formation using selenosulphate has not been unambiguously de-termined. An important difference between depositions using selenosulphate andthiosulphate is that the former are carried out in alkaline solution, in contrast to the(mainly) acidic conditions used for thiosulphate depositions. This means that bothion-by-ion and cluster mechanisms can occur using this reagent, as has beenshown to be the case for CdSe, CdS, and PbSe [15,46]. Once again, it requires em-phasizing that, although selenosulphate depositions are invariably assumed to oc-cur via free selenide ions, this has not been proven, and a complex-decompositionpathway cannot be excluded in selenosulphate depositions in general.

Thermodynamic analyses of metal sulphide formation from thiosulphate[52] and thiourea [53] and metal selenide formation from selenosulphate [13,54]


have been made. These analyses are based on the assumption that free chalco-genide ions are formed.

3.4 NUCLEATION, ADHESION, AND FILM GROWTH

Probably the least-known aspect of the CD process is what determines the nucle-ation on the substrate. Why do adherent films grow under some conditions andpoorly adherent films or even no film at all under others, even when slow precip-itation occurs in solution? In considering this aspect, the two basic mechanisms—hydroxide and ion-by-ion—may behave very differently, although there are alsofeatures in common. When considering nucleation, the anion-mediated mecha-nisms and the complex-decomposition mechanisms will behave similarly in mostcases. Some basic features of nucleation will first be considered, followed by is-sues specific to each.

Film growth can involve continuous nucleation, sticking of colloids fromsolution and growth of individual crystals in the film. The last in particular is con-sidered in Chapter 10, where it is an important factor in the context of quantumsize effects.

3.4.1 General Features of Nucleation and Adhesion

The basic science behind nucleation and forces between materials have beentreated in Chapter 1. For those interested in this section, it is assumed that this ba-sic science is (more or less, at least) understood. However, the basics treated inChapter 1, while important to an understanding of film (as opposed to isolated crys-tal) formation, are not enough by themselves to provide a phenomenological ex-planation of film formation. We would ideally like to be able to predict in advance,from fundamental principles, whether a particular bath formulation will result inadherent films or not. We cannot! However, if we cannot reliably predict adhesion,we can at least choose conditions so that the probability of adhesion is good.

Considering first adsorption of metal ions or neutral species directly on thesubstrate, there are a number of possible mechanisms for this process. Most sim-ply, there will an equilibrium between metal species in solution and a solid sur-face leading to dynamic adsorption of the metal. Adsorption of metal ions ontosolid surfaces has been extensively studied, to a large extent because of the use ofoxide surfaces to adsorb heavy metal ions and remove them from solution (seeRef. 55 for an example and list of other references on this subject). This adsorp-tion may go even farther with ion exchange between the solution metal ions andions in the substrate (again, glass is a good example of where this may occur).

Coulombic attraction between charged species in solution and a surface mayplay a part in initial adsorption on the surface. Under the high pH values more com-


mon in CD, most oxide surfaces (including glass) tend to be negatively charged be-cause of the acid–base equilibrium of the oxide [see Eq. (2.16)]. Positive ions (e.g.,Cd2�, Cd(OH)�) will be attracted to this surface by coulombic forces (in practice,this attraction can be reduced by the solvation shell of the metal ion).

Specific chemical interactions between primary particles or even reactantsand the substrate is another parameter that can aid in prediction of adhesion, atleast qualitatively. Chalcogenide ions will chemisorb to many metals, in somecases forming a surface compound with the metal. Au, Ag, and Cu are the best ex-amples. For Au in a solution of sulphide ions, the surface can be considered a goldsulphide entity; for Cu, bulk sulphidization occurs, leading to eventual disintegra-tion of the entire Cu to CuxS. However, since chalcogenide ion requires time to beformed in a CD process, it is more likely that specific (or even nonspecific, as de-scribed earlier) adsorption of the metal ions or species is dominant in the initialadsorption process. A study of Cd2� adsorption on SiO2 from ethanol/cyclohex-ane mixtures followed by CdS formation by reaction with H2S has concluded thatthe CdS bonds to the silica via SiMOH linkages [56]. While not normally con-sidered, it is possible that the chalcogen presursor is bonded, through the chalco-gen atom, to the surface of some substrates, mainly metals.

Considering now particle adsorption, the section on forces in Chapter 1 con-cluded by stressing that it was normal for particles to stick together (the van derWaals attractive forces eventually dominate in CD processes unless a protectivesurface layer is present—and sometimes even then). Yet this property of “stickingtogether” and also of “sticking to the substrate” clearly can vary greatly; sufficientparticle adhesion to cause aggregation may not be sufficient to form an adherentfilm (adherent meaning that the particles stick both to the substrate and to eachother—both are necessary). However, since colloids do normally aggregate, themain task in CD is to ensure that these aggregated particles (and therefore the pri-mary ones attached to the substrate) do adhere well to the substrate.

Why was the phrase and therefore the primary ones attached to the sub-strate just used? Why not just aggregated particles? It is probably not a very in-accurate generalization to state that adherent films of a reasonable thickness will,in most cases, not form from a solution in which aggregation has already visiblystarted and in which no new product is being formed before the substrate is im-mersed. As a clear example of this, if, during CdSe deposition from a solution inwhich the hydroxide mechanism is operative, the substrate is placed in the depo-sition solution after precipitation starts, although some deposition will occur, it isusually poor quality (poorly adherent and patchy). Clearly the initial adsorptionprocess is important, not just for the initial deposit but for the entire film. This willnot be a surprise to the many who deal with covering a surface—whether by elec-trodeposition, vacuum coatings such as evaporation or sputtering, or, to take evenmore common examples, painting or gluing, where the state of the initial surfaceis very important. However, while a clean substrate is as important for CD as inany coating process, there must be other factors involved.


The fact that primary particles will stick where aggregated ones will not canhave (at least) two causes. One is that the primary particle may be different thanthe final aggregate (e.g., in the hydroxide mechanism, the first stage is adsorptionof hydroxide particles). The other is explained by surface energy: A single parti-cle has a larger external surface-to-volume ratio than an aggregate of the same par-ticles, meaning a larger surface energy and therefore greater potential to stick to asurface in which it comes in contact. While it should always be remembered thatadsorption from solution will certainly reduce this surface energy, perhaps drasti-cally, the difference between single and aggregated particles remains valid.

Finally, and this is likely to be important for nucleation in many cases, theeffect of the electric field (Helmholz layer) at the substrate/solution interface inpromoting formation of a deposit under conditions where none can form in the so-lution (described in Sec. 3.3.2) should be considered. Whether or not this occurs,and to what extent, can be experimentally measured.

3.4.2 Ion-by-Ion Mechanism

Most nucleation studies of CD have treated either the hydroxide or hydroxide-complex mechanisms (see later) or have not clearly defined which mechanismwas, indeed, operative under the conditions of the experiments. Due to the paucityof dependable experimental data, therefore, we consider nucleation and growth bythe ion-by-ion mechanism, to a large extent, from a theoretical viewpoint.

Figure 2.4 showed a general-form curve of film formation as a function oftime. This form is valid in many cases regardless of the mechanism. For the ion-by-ion mechanism, an induction period is generally necessary for sufficientchalcogenide ion to build up and form a solid metal chalcogenide phase. It is prob-able that some metal ions adsorb on the substrate, e.g., by an ion exchange, anelectrostatic mechanism, or simple equilibrium (see Sec. 3.4.1). However, whilethis stage may be important for film initiation, it is not normally considered as agrowth stage in the usual sense of the word. Growth can be considered to beginwhen stable clusters of the deposit begin to form on the substrate.

Ideally, deposition occurs only on the substrate and not in solution. This ispossible due to the effect of a surface (even one in which no chemical interactionoccurs with any constituents of the deposition solution—and such interactionprobably does occur to a greater or lesser extent; water, for example, interacts withmany different surfaces). It is easier to nucleate on a surface than from a homoge-neous solution. The possible effect of the electric field at the substrate/solution in-terface on promoting nucleation has already been described. Additionally, somesurfaces are easier to nucleate on than others. This is the basis of the sensitizationof some surfaces, usually glass by a SnCl2 treatment. The SnCl2 hydrolyzes togive nuclei of hydrated tin oxide on the glass surface, and these nuclei then formnucleation centers for growth of the CD film. Such sensitization may reduce (evento zero) the initial induction (nucleation) time of deposition. An example of this is


the deposition of CdS on a SnO2/glass (conducting glass) surface that has alreadybeen covered with a layer of CdS. Essentially no induction time was found for filmgrowth, compared with a few minutes induction time for deposition, from thesame solution and under otherwise identical conditions, on bare conducting glass[57]. This was explained by catalytic decomposition of the thiourea by the prede-posited CdS, and hence immediate film growth, compared to the slower decom-position of thiourea in the absence of this catalytic surface.

Once nucleation of clusters has begun, growth can occur, since in mostcases the depositing material will deposit on itself more readily than on a substrate(since the reactants, almost by definition, chemically bond to the product). ThusCdS will chemisorb sulphide and/or cadmium ions, depending on the absorptionproperties of the CdS, in particular the crystal face involved. The crystal size ofdeposits formed via the ion-by-ion process tends to be larger than those of thesame material deposited via the hydroxide mechanism (see Sec. 3.4.3). This canbe explained by slower nucleation, resulting in fewer nucleii; growth by homoge-neous formation of chalcogenide ion tends to favor slower nucleation for the ion-by-ion process. These nuclei therefore have more lateral room to grow in the planeof the substrate.* In principle, large crystals (of the order of microns) could be ex-pected in this case. While true in some cases, crystal size was usually muchsmaller than this. Growth termination by adsorption of various species that are notinvolved in the growth process or by defect formation are possible reasons for this.

One possible measure of ion-by-ion growth is that, in contrast to a clustermechanism, deposition of fresh material will generally be preferred on already-ex-isting deposit, as already noted. This means that the number of nucleii may not in-crease greatly (after the very early stages of film formation) and that the film willgrow by growth of these initial crystals. This has been reported, for example, in thecase of Ag2S deposition from a thiourea/ammonia bath [58]. While the mechanismof this deposition is not certain, the combination of this growth pattern togetherwith the strong complexation of the Ag by thiourea (strong AgMS bond in thethiourea complex) suggests an ion-by-ion complex-decomposition mechanism.

Growth of various semiconductors onto certain single-crystal substrates hasresulted in epitaxial growth in a number of cases. This epitaxy has been well stud-ied for CdS deposition by Lincot et al. [59–63]. Although the epitaxy requires acertain degree of lattice matching between semiconductor and substrate, chemicalinteractions between the constituents of the deposition solution and the substrateare important as well (discussed in more detail in Chap. 4). It is a reasonable as-sumption that epitaxial deposition occurs via an ion-by-ion process. Indeed, it has

* It is a characteristic of the ion-by-ion mechanism for CdSe deposition from NTA solutions that thinfilms are highly scattering, and this scattering decreases as film thickness increases. It is likely that thisscattering is due to voids in the thin films, which are a result of the low density of initial nuclei.


been observed that epitaxy ceases when CdS colloids begin to appear in the solu-tion, allowing a cluster growth to occur [63].

Termination of growth due to depletion of reactants will eventually occur.The reactants do not have to be completely used up. The concentration of freemetal ions will decrease not only due to decrease in total metal concentration, butalso because the ratio of the metal ion to complexant (the concentration of the lat-ter usually remains constant throughout the reaction unless ammonia is used in anopen vessel) also decreases. Additionally, as the chalcogen source is used up, thechalcogenide-forming (or chalcogen complexation) reaction slows down. Thus, atsome point, the rate of deposition will slow down to an unacceptable degree, eventhough there may be an appreciable fraction of the reactants remaining in the de-position solution. For the ion-by-ion growth, as the deposition slows and there ismore time for rearrangement of newly formed material to its most stable configu-ration, the likelihood of larger and better-formed crystals may increase, as shownin an early study of PbSe deposition [13].

3.4.3 Cluster Mechanism

The general shape of the growth of CD films as a function of time is often similarfor the cluster mechanism as for the ion-by-ion mechanism (Fig. 2.4). Figure 3.4shows an actual example of CdSe deposition from a solution (containing nitrilo-

FIG. 3.4 Time dependence of CdSe (nitrilotriacetate bath, hydroxide mechanism, roomtemperature) film growth measured by quartz crystal microbalance.


triacetate as complex and under conditions where Cd(OH)2 is present as a colloid)measured by a quartz crystal balance. This deposition was not continued to the in-evitable termination point, but the initial stages of the induction period followedby essentially linear growth are clear.

Since there is a colloidal phase present in the solution from the very begin-ning, why is there an induction period at all in this growth? Should these colloidsnot stick to the substrate and build up a film? The answer, at least in this specificcase (and probably in many others), is yes and no; yes, the colloids will stick tothe substrate, and no, they need not build up to form a film. There are a number ofstudies that show that Cd(OH)2 forms on the substrate from the start of the pro-cess. An early radiochemical study of CdS deposition showed the initial presenceof Cd, free of S, on glass immersed in the (ammoniacal) deposition solution bothwith and without thiourea present [39]. The amount of Cd (considered to beCd(OH)2) was constant with time (after an initial short time) until S began to bedetected in the deposit, at which point it started to increase. Cd(OH)2 was foundon the substrate, using XPS analyses, for CdSe deposition on glass [15] and forCdS deposition on Si using an ammonia bath [37], only using deposition solutionsthat did not contain the chalcogen presursor. The Cd(OH)2 only formed on thesubstrate at a high-enough pH (typically 9 and above, although this value will de-pend greatly on other solution parameters, in particular complex:Cd ratio and tem-perature). In both cases, the coverage of the substrates by the Cd(OH)2 was usu-ally sparse, as evidenced by the predominant substrate (Si) signals, and neverincreased to a level where the Si was not predominant. TEM studies of theCd(OH)2 formation in the CdSe deposition, where selenosulphate was present,confirmed this sparse coverage by Cd(OH)2[15]; CdSe formed only after sometime. Similar results were obtained for PbSe deposition; immersion of Au-coatedglass in an alkaline solution of (selenosulphate-free) citrate-complexed PbAc2 re-sulted in adhesion of hydrated lead oxide (there is no stable simple lead hydrox-ide) to the Au, which was not washed off by water (Fig. 3.5), although in this case,PbSe formed more rapidly than did CdSe when selenosulphate was present [46].

All these results show that Cd(OH)2 colloids do adsorb on a substrate (ei-ther under conditions where Cd(OH)2 is present in solution or, according to thestudies of Rieke and Bentjen and Ortega-Borges and Lincot [48], even when it isnot present in solution but under solution conditions close to solid hydroxide for-mation). The induction period when “no” deposition is seen in the hydroxide-clus-ter deposition therefore is understood to mean that a fast and nongrowingCd(OH)2 adsorption has occurred, which is too fast and/or too little to measure bythe experimental methods used to make the kinetic curves, and that only when thehydroxide starts to convert into the chalcogenide, by reaction of the slowly formedchalcogenide ion with the hydroxide, does real film formation proceed.

An obvious question at this point is: Why does CdS (CdSe) grow, i.e., con-tinue to deposit and form a thick film, but Cd(OH)2 does not? Rieke and Bentjen


discussed the adsorption of various Cd and Cd-OH species onto the Si substrate.The Si will be covered with an oxide (therefore it can be treated similarly to quartzand even approximated to glass for the present purposes), and this oxide will benegatively charged at the pH values involved. Adsorption of positive ions from so-lution (Cd2�, Cd(OH)�), would be favored, and would eventually neutralize thischarge, after which further adsorption of positively charged species would nolonger occur. This scenario, however, was considered to conflict with the obser-vation that greater amounts of Cd(OH)2 were found on the (Si) substrate at higherpH values, where the concentrations of the positive species would be greatly re-duced. Adsorption of neutral Cd(OH)2 was considered more likely. This adsorp-tion could occur in different ways. Direct adsorption of Cd(OH)2 colloids from so-lution was one possibility. However, since Cd(OH)2 could form, at least to someextent, on the substrate from solutions where no Cd(OH)2 was present (see ear-lier), surface-catalyzed adsorption was considered, such as [48].

Cd(NH3)42� � 2OH� � surface site D [Cd(OH)2]ads � 4NH3 (3.56)

where the OH� may or may not be that originally bound at the surface, or, in thehydroxide–chalcogenide complex-decomposition mechanism, by adsorption ofthe ammine-hydroxide species in Eq. (3.52).

The various observations that, at relatively low pH, the coverage of the sub-strate by Cd(OH)2 was poor indicate the dynamic equilibrium between adsorbedCd(OH)2 and the solution; the Cd(OH)2 was in a continual state of dissolution anddeposition. Since the concentrations of both free Cd2� and OH� was constant (in

FIG. 3.5 TEM image of Au film-on-glass immersed for 3 min in a solution of PbAc2 (60mM) and trisodium citrate (160 mM) at pH 10.8. Electron diffraction showed the blackcrystals to be hydrated lead oxide.


the absence of large-scale depletion of the reactants, the case when no chalco-genide precursor was added), this dynamic equilibrium resulted in a steady (andlow) concentration of Cd(OH)2 on the substrate.

When chalcogenide ion begins to form and converts the hydroxide to thechalcogenide (the simple hydroxide mechanism), or when the chalcogenide pre-cursor forms a complex with the surface-adsorbed species followed by complexdissociation (the complex-decomposition mechanism), this disturbs the equilib-rium, allowing more surface-adsorbed hydroxide to form, resulting in filmgrowth. [Note that films of Cd(OH)2 (and other hydroxides) can be grown by CDunder conditions where generation of OH� occurs; i.e., the system is not in equi-librium]. This is one likely scenario for the linear part of the film growth. Thereare other possibilities, and it is likely that more than one mechanism of growth oc-curs. These other possibilities include direct adsorption of colloids from the solu-tion and parallel ion-by-ion growth on the primary deposit.

Direct adsorption of colloids can certainly occur. However, this mechanismis often associated with poorly adherent and optically scattering films. It occurswith these properties if the substrate is placed so that sedimentation of colloidsfrom the solution occurs directly onto the substrate. For this reason, the substrateshould be placed either vertical in the solution or, if non vertical, with only thelower side of the deposit (which forms on both sides of the substrate) retained. Inprinciple, direct adhesion of colloids could result in adherent films, but it is morelikely that this involves isolated colloids or small aggregates that have greater con-tact area than a large aggregate. Greater contact area here means that 200 (for ex-ample) colloids that have adsorbed onto the surface one by one forming a singleaggregate would have a larger contact area to that surface than the same 200 col-loid aggregate that adsorbed, as an aggregate, in a single step. There is also thequestion of whether a Cd(OH)2 colloid would adhere better to a CdS surface thanto itself (for the case where hydroxide is present in solution). Since the hydroxideis in a dynamic equilibrium, nonaggregated particles (or particles containing onlya small number of crystals) will be more likely to be present for the hydroxide thanfor the chalcogenide (where the equivalent equilibrium is likely only for very tinynuclei).

Parallel ion-by-ion growth might occur on previously deposited (by the hy-droxide mechanism) film. However, at least in the simple hydroxide mechanism,where a solid hydroxide is present in solution, this is not likely to be a major fac-tor in the growth, except under solution conditions close to the transition betweenhydroxide and ion-by-ion growth. The reason for this is that the chalcogenide isbeing formed homogeneously throughout the solution. In the case where hydrox-ide is present in solution, most of this chalcogenide will react with colloidal metalhydroxide in solution (as seen in most cases by precipitated metal chalcogenide),and the concentration at the substrate will be very low—enough to convert hy-droxide to chalcogenide perhaps, but less likely to form a new phase.


For the case where no hydroxide is present in solution but is formed on thesubstrate while the homogeneously formed chalcogenide will not react with hy-droxide in solution, some or even much of it may still not reach the substrate. Thisis due to the relative lack of stability of the chalcogenide ions, which could be ox-idized homogeneously, e.g., by dissolved oxygen. This will be more important forselenide than for the considerably more stable sulphide; but even sulphide is notstable in very low concentrations unless oxygen is rigorously excluded. This po-tential problem will not exist for the complex-decomposition mechanism, wherethere is a high concentration of the chalcogenide complex.

The crystal size of the deposit obtained by the simple hydroxide mechanismwill be essentially that of the Cd(OH)2 clusters, which are converted into, e.g., CdS.Such colloids tend to be very small (typically in the region of 5 nm), and thereforethe CdS crystal size should be similar. It may be larger if some form of ion-by-iongrowth occurs on these primary crystals. Also, it may be different if the hydroxideis present on the substrate but not in solution, since the size of the hydroxide depositon the substrate will, at least in part, be affected by different factors than that formedhomogeneously in the solution. In general, crystal size in the simple hydroxidemechanism (hydroxide colloid present in solution) is smaller than that formed in anion-by-ion process. This has been shown for CdS, CdSe, and PbSe, where typicalcrystal sizes for the hydroxide (ion-by-ion) mechanisms were found to be (in nm):CdS—5 (�70); CdSe—5 (15); PbSe—5 (10–1000) [15]. In contrast to the ion-by-ion mechanism, the crystal size of films formed via a cluster mechanism is not ex-pected to grow greatly (some modest growth can and usually does occur) with in-creasing film thickness, since film growth occurs by sequential addition of newclusters. Thus, in our observation of CdSe growth via the hydroxide mechanism,the color of the film (a sensitive measure of crystal size due to size quantization; seeChap. 10) changes only a little during the deposition (color refers to spectral posi-tion and not to depth of color, which, of course, does increase with film thickness).

SEM micrographs in the study of Rieke and Bentjen [37] showed that, al-though the number density of CdS nucleii/unit area on the substrate was constantwith time after deposition started, the size of the nucleii increased linearly withtime. Additionally, the size distribution of the nucleii, both in the early and laterstages of growth, was quite narrow. These film growth kinetics were identifiedwith burst-type nucleation, well known in homogeneous solution precipitation,where a homogeneous reaction in solution causes sudden nucleation whenever acritical concentration of one of the reactants is reached. This nucleation reducesthe concentration of this reactant so that further growth occurs only on existing nu-cleii (nucleation usually requires a large supersaturation, while growth on an ex-isting nucleus does not). This type of nucleation usually results in a narrow sizedistribution, as seen here.

In a quartz crystal microbalance investigation of CdSe film growth rate froma selenosulphate/ammonia/triethanolamine bath with different Cd:selenosulphate


ratios, two peaks were observed in the growth rate vs. time plot of all solutions[64]. This was explained by a two-stage growth. From electron microscopic ex-amination of the growing films, the first stage was attributed to instantaneous nu-cleation and 2-D lateral growth to cover the substrate, while the second stage wasdue to 3-D nucleation and growth at random sites on the first layer. Aggregationof colloidal particles was invoked as the mechanistic pathway.

After all the speculation involved in the foregoing discussion of filmgrowth, the termination step is as simple to explain as for the ion-by-ion mecha-nism. Growth can occur as long as the concentration of chalcogenide anion is highenough to allow Reaction (3.5), the conversion of the hydroxide to the chalco-genide, to occur [or, for the complex-decomposition mechanism, sufficientchalcogenide precursor as, e.g., in Eq. (3.7)]. It is also possible that depletion ofthe metal hydroxide occurs first (in which case the mechanism may change to ion-by-ion, as described earlier). In either case, termination is simply due to depletionof the reactants. Typically in the cluster mechanism, most of the reactants are lostin homogeneous precipitation.

3.4.4 Complex-Decomposition MechanismThe initial nucleation stage of the complex-decomposition mechanism is probablysimilar to the simple free-anion mechanism. Either ionic or molecular metalspecies (ion-by-ion) or Cd(OH)2 (cluster) adsorbs on the substrate. However, in-stead of conversion of the hydroxide to sulphide by topotactic reaction with sul-phide ions, the chalcogenide precursor (in almost all studies of this mechanism,that is thiourea) adsorbs on the Cd(OH)2 surface to form a hydroxide–thioureacomplex, which then decomposes to CdS.

A possible difference between the simple hydroxide mechanism and thecomplex decomposition is in the manner of crystal growth. We noted earlier thatcrystal growth in the simple hydroxide mechanism may occur via an ion-by-ionprocess but probably not to a large degree, since the concentration of chalcogenideion would be very low indeed (it would react rapidly with Cd(OH)2 in the solu-tion). However, for the hydroxide-decomposition mechanism, the chalcogenidereactant is not the free chalcogenide ion but the precursor, which is present inmuch higher concentrations. Therefore once a solid phase capable of catalyzingthe chalcogenide precursor has formed, the crystal growth is quite likely toswitch over to a predominantly ion-by-ion process, such as in Reaction (3.51).This means that, as for the pure ion-by-ion process, the crystal size might beexpected to be larger than for the pure hydroxide mechanism, since the ion-by-iongrowth favors, in principle, crystal growth rather than renucleation. Typicalcrystal sizes for CdS prepared from the ammonia/thiourea bath, which appears un-der many experimental conditions to proceed via the hydroxide-complex mecha-nism, are in the region of 10 nm to several tens of nanometers, larger than the ap-proximately 5 nm obtained for CdS and CdSe from the simple hydroxide-clustermechanism.


3.5 KINETICS OF DEPOSITION

Kinetic studies have been a popular topic in CD, particularly for CdS. This wouldsuggest that the present section will be a large one to reflect this activity. In fact,the reverse will be the case. This will be a relatively short section that will not tryto cover even a moderately large part of the kinetic studies. The reason is that ki-netic measurements have been used, to a large extent, to study the mechanisms ofdeposition, and this has been dealt with already (not the details of the kinetic stud-ies, but the conclusions). Additionally, since a CD process can vary widely inrate—from a few minutes to days and weeks—often depending strongly on smallchanges in one or another concentration of a particular reactant, the important in-formation to be learned from kinetic studies (apart from mechanistic diagnosis) ishow this rate depends on experimental conditions, and this can be done with a fewselected examples from the literature.

Ortega-Borges and Lincot [48] carried out a detailed kinetic study of CdSdeposition from the standard ammonia (ammonium)/thiourea bath using a quartzcrystal microbalance to measure film thickness. They measured a deposition ratewith fractional values of reaction order

rate � K (3.57)

They therefore concluded that several different rate-determining steps were in-volved in the deposition. Figure 3.6 shows the dependence of the deposition rateon the concentration of the reactants (Cd, thiourea, ammonia, and pH—the lastvaried through introduction of ammonium ion) (a) as well as an Arrhenius plot ofthe deposition (b) for the CdS deposition. From the kinetic data, they deduced thehydroxide-complex-decomposition mechanism, given earlier in Eqs. (3.50) and(3.51) and, more specifically, as

Cd(NH3)42� � 2OH� � surface site D Cd(OH)2(ads) � 4NH3 (3.58)

This first step represents a reversible adsorption of Cd(OH)2 on the substrate. Thereaction order of 1.5 for hydroxide [given by H� in Eq. (3.57)] implied the par-ticipation of two hydroxide ions in the process. The next step was formation of acomplex between the adsorbed Cd(OH)2 and thiourea:

Cd(OH)2(ads) � (NH2)2CS → Cd(OH)2�SC(NH2)2(ads) (3.59)

This complex then decomposes to CdS:

Cd(OH)2�SC(NH2)2(ads) → CdS � CN2H2 � 2H2O (3.60)

probably by nucleophilic attack of the thiourea S atom on the Cd(OH)2. The acti-vation energy of the deposition (85 KJ/mole), measured from the Arrhenius plotin Figure 3.6b, is very similar to that of thiourea decomposition, suggesting theslow nature of Reaction (3.60). Earlier measurements of PbSe deposition from a

[Cd]0.6[Tu]0.8

��[NH3]3.3[H�]1.5


selenourea bath gave an activation energy of 60 KJ/mole [65], with the lowervalue compared to CdS deposition presumably reflecting the greater instability ofselenourea compared to thiourea.

Using the same basic system and similar experiments, Doña and Herreromeasured reaction orders for the various species comparable to those measured inthe study of Ortega-Borges and Lincot, except for that of ammonia, which was 1.8instead of 3.3 [49]:

rate � K (3.61)

This, together with the known tendency of metal ions to form mixed hydroxy–am-mine complexes, suggested to them that two ammonia molecules were involvedin the first step and that the adsorbed species in Reaction (3.58) was a hy-droxy–ammine species, viz. Cd(OH)2(NH3)2. Decomposition of the hydroxide–ammine–thiourea complex was then assumed to occur by nucleophilic attack of anammonia species on the SBC bond of the thiourea.

Based on this mechanism, a detailed theoretical model has recently beenproposed for CdS deposition from the thiourea/ammonia bath [65a]. Prediction ofdifferent aspects of the deposition kinetics using this model provided a very goodfit with the relevant experimental data.

Rieke and Bentjen [37] studied the kinetics of CdS deposition from an am-moniacal thiourea bath using SEM. As discussed earlier, they found that good-

[Cd]0.9[Tu]1.1[OH�]1.7

��[NH3]1.8

FIG. 3.6 (a) Log–log plots of reactant concentration (X) vs. deposition rate for CdS de-posited from an ammonia/thiourea bath. Standard conditions: [Cd] � 14 mM; [Tu](thiourea) � 28 mM; [NH3] � 1.74 M; T � 60°C. Reactants are Cd (total concentration);Tu; NH3; pH (adjusted by adding ammonium ion) that gives hydroxide concentration.(b) Temperature dependence of deposition rate. (Adapted from Ref. 48).


quality films formed in a pH range where Cd(OH)2 formed on the substrate (Si, intheir experiments) but not in the bulk of the solution. Their kinetic study was madein this pH range (specifically, pH � 9.55). No CdS deposited initially, but the rateof formation of CdS increased with time, eventually becoming more or less con-stant over the time of their experiment. This is characteristic of an autocatalytic re-action, where the initial deposit accelerates the rate of further deposition.

O’Brien and Saeed, using ethylenediamine as compexant, higher depositiontemperatures, and glass as a substrate, found that the thickness of the CdS film in-creased linearly with time (after an initial induction period) and also that there wasno increase in the size of the nucleii (both in contrast to the previous study) [40].In spite of the different experimental conditions, the mechanism of the depositionsin both studies appears to be essentially the same, i.e., hydroxide-mediated catal-ysis of thiourea decomposition.

It must be kept in mind that the kinetics of CD, as with the deposition mech-anism, can be very different from one system to another. Two connected exam-ples of this are given here.

In one study of PbSe deposition from a citrate-complexed selenourea solu-tion containing hydrazine, the rate was proportional to the pH and to the sele-nourea concentrations but independent of the Pb and citrate concentrations [65].This was explained by a rate-determining step involving decomposition of sele-nourea at the (catalytic) PbSe surface by hydroxide. It is noteworthy that the Pbconcentration was typically an order of magnitude less than that of selenourea.Therefore the independence of the rate on Pb (or citrate, which determines theconcentration of free Pb2�) concentration, suggests that formation of selenide ion,and not a complex-decomposition mechanism, occurs.

The second example is seen in the study of PbSe deposition by Kainthla etal. from selenosulphate solution [41]. In most examples of CD from alkaline so-lution, the deposition rate increases with increase in pH. This is due to both thegreater rate of decomposition of the chalcogenide precursor at higher pH (this de-composition usually involves hydroxide ions) and, in many cases, the greaterprobability of solid hydroxide formation (as long as this is not excessive). How-ever, for PbSe deposition using citrate as complex for the Pb and selenosulphateas Se precursor, the opposite occurs: The deposition rate decreases with increasein pH. This is due to the specific hydroxy–citrate complex formed:

Pb(OH)C6H5O72�D Pb2� � OH� � C6H5O7

3� (3.62)

Increase in pH (� increase in [OH�]) shifts the equilibrium to the left, resultingin a lower concentration of free Pb2� ions and thus a slower reaction to give PbSe.This means that, in contrast to the deposition from a selenourea bath described ear-lier, the rate is dependent on Pb concentration and possibly independent of hy-droxide concentration at a constant free-Pb2� concentration. This would then sug-gest that the opposite mechanism, i.e., a complex decomposition, is effective forthe selenosulphate bath. It is stressed that these conclusions on selenide formation


or complex-decomposition mechanisms are indications of which mechanism istaking place, but are far from being firm proof of this.

For the deposition of ZnS and ZnSe, hydrazine is normally used to formfilms at a reasonable rate. The role of the hydrazine is not obvious. It is temptingto assume that hydrazine, being a strong reductant, reduces the chalcogen precur-sor to chalcogenide ion, as was assumed for CdTe deposition (see Sec. 3.2.3).However, this appears to be oversimplified. In a study of the effect of variousamines (including hydrazine) on the deposition rate (and composition) of ZnSfilms deposited from ammonia/thiourea baths, Mokili et al. found a strong depen-dence of the rate on the type of amine added [66], as shown in Figure. 3.7. Whileit is difficult to separate the effect of concentration from the different types ofamines in this experiment, it is clear that an increase in rate is general on additionof amine (apart from the initial induction time using triethanolamine). Since theamines also act as complexation agents, they would, on this basis, be expected toreduce the deposition rate (by reducing the free Cd2� concentration). The fact thatthe opposite occurs implies that they must increase the thiourea (or selenosulphatefor selenides) decomposition. In this respect, all the amines used have pronouncedreducing properties, with redox potentials of �0.46 (triethanolamine), �0.56(ethanolamine), and �1.16 (hydrazine), which parallels the order of increasein deposition rate seen in Figure 3.7. O’Brien et al. discuss various theories for

FIG. 3.7 Evolution of ZnS film thickness with time. ZnCl2/NH4Cl/thiourea/NH3 bath at85°C. The effect of various amines on growth. Triethanolamine—0.2 M; ethanolamine—0.7 M; hydrazine—3 M. (Adapted from Ref. 66).


the accelerating effect of hydrazine on the CD process [67]. While agreeingthat the hydrazine increases thiourea decomposition, the specific details of theeffect of hydrazine (and other amines) on the CD process are still not fullyunderstood.

The activation energies of the deposition for both ZnS [68] and ZnSe [69],measured from Arrhenius plots, are 21 and 26 KJ/mole, respectively, muchsmaller than the values for CdS (85 KJ/mole) or PbSe (60 KJ/mole) described ear-lier. Stirring does not affect the deposition rate for either ZnS or ZnSe, so the de-position is not under diffusion control. In interpreting activation energies for CDprocesses, it is important to remember that what is measured is the film growth,and this is not necessarily the same as the rate of formation of the metal chalco-genide, much of which is usually formed homogeneously in the solution. The Zncompounds were both probably formed by a cluster mechanism, in contrast to theion-by-ion complex-decomposition mechanism probable for the CdS. The “acti-vation energy of deposition” for the ZnS(Se) therefore depends to a large extenton the rate of sticking of clusters, although other factors could also be involved,such as a parallel ion-by-ion (whether by complex decomposition or free chalco-genide ions). To interpret such results correctly will require a study of activationenergies of different compounds deposited under different and controlled mecha-nistic pathways and preferably also measuring the total amount of product formed(in solution as well as on the substrate and other surfaces). It is also relevant thatthe crystal size for ZnS and ZnSe deposits is typically smaller than for CdS orCdSe deposited from an ion-by-ion bath; this supports a cluster mechanism forthese depositions. In fact, in contrast to CdS and CdSe, there are no cases in theliterature where ZnS or ZnSe have been clearly shown to have been deposited viaan ion-by-ion mechanism.

3.6 DEPOSITION FROM ACIDIC BATHS

While the majority of CD reactions have been carried out in alkaline baths, therehave been a number of sulphide depositions reported in acidic baths. These are all(with one exception, Sn-S from acidic nonaqueous S baths [70]; see Chap. 6)based on two sulphide precursors: thiosulphate (more commonly used in the veryearly studies, but still sometimes used) and thioacetamide. Beutel and Kutzelnigg[71] described a large selection of colored films deposited on various metals us-ing CD (and also electrodeposition) from metal salt–thiosulphate solutions. Nocharacterization of these films (other than interference colors) was made. How-ever, it is clear that CD of metal sulphides did occur in some of these cases (seeChap. 6). Lokhande described a variety of sulphides (CdS, ZnS, Bi2S3, Sb2S3,As2S3, Cu2S, Ag2S, SnS2, and PdS2) deposited using thiosulphate at a pH of typ-ically ca. 3 [72]. Other studies on deposition from acidic thiosulphate baths are:PbS [73], Ag2S [74], Bi2S3 [75], CuxS [76,77], Sb2S3 [78], SnS2 [79].


The mechanism of the thiosulphate reaction is not clear. Lokhande, in hisstudies, has suggested an internal reduction involving the reaction

2S2O32� → S4O6

2� � 2e� (3.63)

These electrons reduce elemental S formed in Reaction (3.21) or (3.22) to givesulphide (or hydrosulphide) ions, as in Reaction (3.24), which react with the metalions. The general mechanism of sulphide generation has been assumed in moststudies using thiosulphate, e.g., Nair et al [78] for Sb2S3 deposition and Groz-danov et al. in their study of CuxS deposition from Cu-thiosulphate solutions [76],although in this latter study it is noted that the mechanism may be more compli-cated in this Cu-S system.

As has already been pointed out, in spite of the fact that the free sulphidemechanism is invariably assumed in thiosulphate depositions, there is no evidenceup to now against the complex-decomposition mechanism. No thorough mecha-nistic or kinetic studies have been made on this system. Since the studies on CDusing thiosulphate have not attempted to differentiate between these differentmechanisms (and since such differentiation may be difficult, this is not surpris-ing), we are left with the conclusion that there is no clear consensus on whichmechanism is operative. Also, the mechanism may vary depending on conditions(as for the alkaline baths), and of course a combination of mechanisms may be op-erative in some cases.

Other points to note when considering thiosulphate as a reagent in CD is thatthiosulphate is a strong complex for a number of metals and, since it is a fair re-ducing agent, also may reduce the metal ions (as is known to occur for Cu2�

to Cu�).Many sulphides have been deposited using thioacetamide in acidic solutions

(Chapter 6 describes most of these). For depositions using thioacetamide, as withthiosulphate, there are no detailed mechanistic studies. Both H2S formation andcomplex decomposition are possible in acid solutions, as discussed in Section3.2.1.3. Deposition of CdS was accomplished using thioacetamide in acidic solu-tion by exploiting electrolytic proton reduction to increase the pH locally at thecathode (substrate), and the mechanism was believed to be a surface-catalyzed de-composition of a Cd–thioacetamide complex [80].

Because of the acidic conditions, with the exception of very acidic cations,nucleation resulting from a solid colloidal phase is unlikely. For a deposition thatoccurs through free (hydro)sulphide ion, it is probable that nucleation occurswhen the concentration of this sulphide ion is high enough to cause precipitationof the metal sulphide, possibly catalyzed by the substrate surface. Similarly for acomplex-decomposition pathway, a high-enough concentration of the final prod-uct to permit solid-phase formation will be required. It should be remembered thatthe concentration of free metal ion will, in most cases, be higher than for alkalinebaths, due to the (usually) weaker complexing strengths of the sulphide precur-


sors. This will allow formation of a solid phase at lower concentrations of sulphidethan in more heavily complexed alkaline solutions.

There appear to be no cases of selenide deposition from acidic baths. Se-lenosulphate is not stable under even mildly acidic conditions, and all selenourea-based baths have been alkaline ones.

Oxides or hydroxides have been deposited from acid baths, in particularreadily hydrolyzable acidic metal ions. These have been discussed in Section3.2.4.

3.7 EFFECT OF STIRRING

There have been only a few studies on the effects of stirring the deposition solu-tion on the deposited film. Overall, stirring affects CD films mainly by preventingdeposition of loosely adhering, large aggregates. These loose deposits are readilyremoved by the stirring action. This is important, since they block the substrate,preventing normal adherent film growth. Such nonadherent deposits can also beprevented without stirring by placing the substrate in the bath at an angle; the de-posit on the upper surface, which will usually be a mixture of adherent and looselyadherent material, can be removed (by wiping with a reagent that dissolves thefilm, often dilute HCl), leaving the film on the lower surface, which does not col-lect precipitated deposit.

Such loosely adhering CdS films in nonstirred solutions have been reportedby Kaur et al. [34] and by Doña and Herrero [49]. The latter and also Ortega-Borges and Lincot [48] found that the rate of deposition is affected by stirring onlyat low stirring rates, and the effect is not large. There is no apparent difference be-tween low and fast stirring rates. This implies that even slow stirring is enough toprevent sticking of large, loosely adhering particles. Apart from deposition rate,the study by Kaur et al. found that stirring could, in some cases, strongly affect thefilm quality and that the effect of stirring was dependent on the concentration ofammonia (relative to the Cd). For low ammonia concentrations, where a visibleCd(OH)2 phase existed, and for high ammonia concentrations, where it is proba-ble that no solid hydroxide phase occurred, strongly stirred solutions resulted information of CdS films with predominantly wurtzite structure, while films de-posited from unstirred solutions contained a large amount of zincblende CdS. Pre-cipitate formed homogeneously in the solution was found to be zincblende, andthe effect of stirring was to minimize sticking of colloidal particles from solution,thereby reducing the zincblende component in the films (and also resulting inmore adherent and specularly reflecting films—adsorption of large aggregatedcolloids from solution caused light scattering and reduction in specular re-flectance). For the intermediate case where there was just enough ammonia to dis-solve the visible Cd(OH)2 (but a colloidal Cd(OH)2 probably was present in solu-tion), stirring did not have any major effect—thick, powdery, and zincblende


films were obtained in both cases. This is unexpected, since these are conditionswhere good films are often obtained.

There is one unusual case where stirring had a very strong effect on the de-position rate: Ag2S deposited from a thiourea bath [81]. The deposition was veryslow in the unstirred solution and increased, more rapidly at first, then linearly, upto the maximum stirring speed of ca. 1100 rpm. The energy of activation of thisdeposition was 20.4 kJ/mole, much less than values typical of reaction-controlleddepositions under similar conditions involving, e.g., CdS or PbS, and similar tothose found for ZnS and ZnSe. However, the Zn-S(Se) depositions were indepen-dent of stirring. It therefore appears that the mechanism of this Ag2S deposition isdiffusion controlled and is unlike other mechanisms discussed previously.

As pointed out by Ortega-Borges and Lincot, the relative independenceof the majority of CD processes on hydrodynamic conditions explains theexcellent lateral homogeneity characteristic of this technique, since the depositiondepends not so much on mass transport in the solution as on chemical reactionrates.

REFERENCES1. WHR Shaw, DG Walker. J. Am. Chem. Soc. 78:5769, 1956.2. DF Bowersox, DM Smith, EH Swift. Talanta 3:282, 1960.3. OM Peeters, CJ de Ranter. J. Chem. Soc. Perkin II 1832, 1974.4. R Williams, PN Yocom, FS Stofko. J. Colloid Interface Sci. 106:388, 1985.5. WH Perkins, AT King. J. Chem. Soc. 103:301, 1913.6. WN Arnquist. Proc. IRE 47:1420, 1959.7. CJ Milner, BN Watts. Nature 163:322, 1949.8. GA Kitaev, TP Sokolova. Russ. J. Inorg. Chem. 15:167, 1970.9. GA Kitaev, VM Makova. USSR Patent 377,445:C.A. 79:84806j, 1973.

10. RA Zingaro, FC Bennett Jr, GW Hammar. J. Org. Chem. 18:292, 1953.11. RA Zingaro, DO Skovlin. J. Electrochem. Soc. 111:42, 1964.12. AA Velykanov, EK Ostrovskaya, NP Garina, VA Turacova, AA Tchurkan. Ukr.

Chim. Zh. 49:764, 1983.13. GM Fofanov, GA Kitaev. Russ. J. Inorg. Chem. 14:322, 1969.14. GA Kitaev, TS Terekhova. Russ. J. Inorg. Chem. 15:25, 1970.15. S Gorer, G Hodes. J. Phys. Chem. 98:5338, 1994.16. GA Kitaev, GM Fofanov. Zh. Prikl. Khim. 43:1694, 1970.17. GK Padam, SK Gupta. Appl. Phys. Lett. 53:865, 1988.18. RW Buckley. In: 11th ECPV Solar Energy Conf: Montreux, Switzerland, 1992,

p 962.19. AV Kalyakina, RI Pelyukpashidi. Tr. Khim. Met. Inst. Akad. Nauk. Kaz. SSR

17:114, 1973.20. VB Patil, PD More, DS Sutrave, GS Shahane, RN Mulik, LP Deshmukh. Mater.

Chem. Phys. 65:282, 2000.21. M Sotelo-Lerma, RA Zingaro, SJ Castillo. J. Organomet. Chem. 623:81, 2001.


22. L Gordon. Anal. Chem. 24:459, 1952.23. HH Willard, L Gordon. Anal. Chem. 25:170, 1953.24. M Izaki, T Omi. J. Electrochem. Soc. 144:L3, 1997.25. M Izaki, O Shinoura. Electrochem. Solid State Lett. 1:215, 1998.25a. M Izaki. O Shinoura. Adv. Mater. 13:142, 2001.26. W Mindt. J. Electrochem. Soc. 117:615, 1970.27. TP Niesen, MR DeGuire. J. Electroceram. 6:169, 200128. DH Klein, L Gordon, TH Walnut. Talanta 3:177, 1959.29. L Gordon, JI Peterson, BP Burtt. Anal. Chem. 27:1770, 1955.30. VK LaMer, RH Dinegar. J. Am. Chem. Soc. 73:380, 1951.31. L Gordon, ML Salutsky, HH Willard. Precipitation from Homogeneous Solutions.

New York: Wiley, 1959.32. JR Gump, GR Sherwood. Anal. Chem. 22:496, 1950.33. PFS Cartwright. Analyst 92:319, 1967.34. I Kaur, DK Pandya, KL Chopra. J. Electrochem. Soc. 127:943, 1980.35. ML Breen, JT Woodward, DK Schwartz, AW Apblett. Chem. Mat. 10:710, 1998.36. GA Kitaev, AA Uritskaya, SG Mokrushin. Russ. J. Phys. Chem. 39:1101, 1965.37. PC Rieke, SB Bentjen. Chem. Mat. 5:43, 1993.38. RO James, TW Healy. J. Colloid Interface Sci. 40:53, 1972.39. ND Betenekov, VP Medvedev, AS Zhukovskaya, GA Kitaev. Sov. Radiochem.

20:524, 1979.40. P O’Brien, T Saeed. J. Crystal Growth 158:497, 1996.41. RC Kainthla, DK Pandya, KL Chopra. J. Electrochem. Soc. 127:277, 1980.42. H Cachet, M Froment, G Maurin. J. Electroanal. Chem. 406:239, 1996.43. P O’Brien, J McAleese. J. Mater. Chem. 8:2309, 1998.44. D Lincot, RO Borges. J. Electrochem. Soc. 139:1880, 1992.45. S Gorer, A Albu-Yaron, G Hodes. J. Phys. Chem. 99:16442, 1995.46. S Gorer, A Albu-Yaron, G Hodes. Chem. Mater. 7:1243, 1995.47. MK Norr. J. Phys. Chem. 65:1278, 1961.48. R Ortega-Borges, D Lincot. J. Electrochem. Soc. 140:3464, 1993.49. JM Doña, J Herrero. J. Electrochem. Soc. 144:4081, 1997.50. EB Andersen. Z. Phys. Chem. B. 32:237, 1936.51. E Beutel, A Kutzelnigg. Z. Elektrochem. 36:523, 1930.52. GA Kitaev, AA Uritskaya. Russ. J. Appl. Chem. 72:592, 1999.53. LE Yatlova, AA Uritskaya, GA Kitaev, TI Dzyuba. Russ. J. Appl. Chem. 66:1534,

1993.54. GA Kitaev, AZ Khvorenkova. Russ. J. Appl. Chem. 71:1325, 1998.55. J Liu, SM Howard, KN Han Langmuir. 9:3635, 1993.56. I Dekany, L Turi, G Galbacs, JH Fendler. J. Colloid Interface Sci. 195:307, 1997.57. D Lincot, J Vedel. In: 10th ECPV Solar Energy Conf. Lisbon Portugal: 1991, p 931.58. TP Bol’shchikovsa, GA Kitaev, VI Dvoinin, MV Degtyarev, LM Dvoskina. Izv.

Akad. SSSR, Neorg. Mater. 16:387, 1980.59. D Lincot, R Ortega-Borges, M Froment. Appl. Phys. Lett. 64:569, 1994.60. M Froment, MC Bernard, R Cortes, B Mokili, D Lincot. J. Electrochem. Soc.

142:2642, 1995.61. R Cortes, M Froment, B Mokili, D Lincot. Philos. Mag. Letts. 73:209, 1996.


62. D Lincot, B Mokili, R Cortes, M Froment. Microsc. Microanal. Microstruct. 7:217,1996.

63. MJ Furlong, M Froment, MC Bernard, R Cortés, AN Tiwari, M Krejci, H Zogg, DLincot. J. Crystal Growth 193:114, 1998.

64. H Cachet, H Essaaidi, M Froment, G Maurin. J. Electroanal. Chem. 396:175, 1995.65. AB Lundin, GA Kitaev. Inorg. Mater. 1:1905, 1965.65a. M. Kostoglou, N. Andritsos, AJ Karabelas. Ind. Eng. Chem. Res. 39:3272, 2000.66. B Mokili, M Froment, D Lincot. J. de Phys. IV 5:261, 1995.67. P O’Brien, DJ Otway, D Smith-Boyle. Thin Solid Films 361:17, 2000.68. JM Doña, J Herrero. J. Electrochem. Soc. 141:205, 1994.69. JM Doña, J Herrero. J. Electrochem. Soc. 142:764, 1995.70. RD Engelken, HE McCloud, C Lee, M Slayton, H Ghoreishi. J. Electrochem. Soc.

134:2696, 1987.71. E Beutel, A Kutzelnigg. Monats. 58:295, 1931.72. CD Lokhande. Mater. Chem. Phys. 28:145, 1991.73. KM Gadave, SA Jodgudri, CD Lokhande. Thin Solid Films 245:7, 1994.74. SS Dhumure, CD Lokhande. Thin Solid Films 240:1, 1994.75. CD Lokhande, AU Ubale, PS Patil. Thin Solid Films 302:1, 1997.76. I Grozdanov, CK Barlingay, SK Dey. Thin Solid Films 250:67, 1994.77. I Grozdanov, M Najdoski. J. Solid State Chem. 114:469, 1995.78. MTS Nair, Y Pena, J Campos, VM García, PK Nair, J. Electrochem. Soc. 145:2113,

1998.79. CD Lokhande. J. Phys. D: Appl. Phys. 23:1703, 1990.80. K Yamaguchi, T Yoshida, T Sugiura, H Minoura. J. Phys. Chem. B 102:9677, 1998.81. H Meherzi-Maghraoui, P Cowache, D Lincot, M Dachraoui. J. Chim. Phys. 96:259,

1999.


4II–VI Semiconductors

There has been a clear emphasis, in the CD literature, on II–VI semiconductors,mostly CdS, some CdSe, and recently on ZnS. This being the case, the reader mayreasonably expect this chapter to be a voluminous one. On the other hand, manyof these studies have focused on deposition mechanisms and kinetics (which aredealt with in the previous chapter), with photovoltaic cells, and, to a lesser extent,with quantum size effects, both of which will be dealt with in subsequent chapters.Two detailed descriptions of the experimental procedure (for CdS and CdSe) aregiven in Chapter 2. This leaves the obvious question: “What’s left?” The presentchapter will answer that question. This includes properties of the films not ex-plicitly discussed in other sections, such as crystal structure, optical and electricalproperties, as well as variants of the deposition process. Also, more detail will begiven on non-Cd chalcogenides. In short, there is indeed much left.

4.1 CdS

A point concerning CdS deposition. Many studies have used what is referred hereto as the standard deposition bath. This bath is made up of a Cd salt, ammonia(sometimes with an ammonium salt to lower the pH) to complex the Cd and ad-just the pH, thiourea, and deposition temperatures usually in the range of60–90°C. Of course, this bath still allows for large differences in reactant concen-trations (e.g., the Cd concentration varies from a low of 1 mM to as much as 100


mM). If the reader has read the previous chapter, it should be obvious that not onlyis the concentration of various reactants important, but so is the ratio between theCd (or other metal ion) and the complexant. With this caveat, we will use the termstandard bath to cover all concentrations, unless there is a specific reason to dootherwise. Since the majority of studies on CdS used this standard bath in oneform or another, the films discussed in this section can be assumed to have beendeposited from such a bath unless otherwise stated. Also for this reason, it is morenatural to begin with properties of the films and afterwards to discuss variationsin deposition.

4.1.1 CrystallographyIn several cases where epitaxial growth occurs involving the ion-by-ion mecha-nism, the crystal structure is dictated by the substrate structure. This is treated sep-arately in Section 4.1.5.

Many papers state that one or the other crystal modification is obtained,without giving either diffraction data or/and where the data is ambiguous. The en-ergy difference between the hexagonal (wurtzite) and cubic (zincblende) phases isvery small (the former is slightly more stable); hence both are often found to-gether. This commonly leads to the presence of twins and stacking faults in thecrystals. The density of stacking faults in films deposited from a standard bath in-creases with increased thiourea concentration or decreased Cd concentration, andis typically 1011–1012 cm�3 [1]. If the cubic phase is annealed at ca. 400°C orabove, the hexagonal phase is normally obtained. In view of the lack of a con-vincing explanation of why one or the other crystal structure is formed, a samplingof reported crystal structures is given in list form, together with differences inpreparation from the standard bath (if any) that may give clues to the crystal mod-ification obtained.

4.1.1.1 Hexagonal (Wurtzite)

When insufficient NH3 was added and Cd(OH)2 was present as a clearly vis-ible suspension, hexagonal CdS formed on the substrate if the solutionwas well stirred; i.e., the precipitate in solution, which was cubic, was notallowed to accumulate on the substrate [2].

A high resolution transmission electron microscopy (HTEM) study of theearly stages of CdS deposition on a carbon-coated TEM grid showed onlyhexagonal CdS to be formed, while hexagonal with some cubic CdS wasformed by precipitation in the solution [3].

Using CdI2 in a standard bath, hexagonal CdS was obtained. (If CdCl2 wasused, the deposit appeared to be zincblende, although it may also havebeen highly textured wurtzite) [4].)

Using CdCl2 (Cd(Ac)2 gave no XRD) hexagonal CdS with moderate texture(0002) was deposited on glass [5].


With ethylenediamine as complexant and with Cd(OH)2 present in solution,some hexagonal CdS was formed, although cubic CdS might also havebeen present [6].

A citrate/ammonia bath gave predominantly hexagonal CdS [7].Some particularly clear examples of predominantly hexagonal formation

were from an acid bath using thioacetamide [8], a triethanolamine bathunder conditions where ion-by-ion deposition was believed to occur andthe deposition rate was slow [9], and a nitrilotriacetate bath, where depo-sition was also slow but a hydroxide cluster deposition was shown to takeplace [10]. (Ion-by-ion growth, under conditions similar to the last exam-ple, but with a high enough complex concentration to prevent Cd(OH)2

formation, also showed apparently highly textured hexagonal CdS, al-though in this case the predominant presence of this form was not unam-biguous [10].)

4.1.1.2 Cubic (zincblende)

If just enough ammonia was added to dissolve Cd(OH)2, the cubic form wasobtained [2] (regardless of stirring—see Ref. 2).

From a standard solution although the concentrations of reactants used werenot given [11].

Mainly cubic obtained from a standard bath on SnO2/glass over a range ofconditions (including with and without ammonium buffer and usingethylenediamine instead of ammonia) [12].

From a bath with low Cd concentration (1 mM) and high ammonia concen-trations (2 M) suggesting that the conditions were such as to favor ion-by-ion deposition [13]. Another study with low Cd concentration (2–5 mM),ammonia concentrations ca. 300 times higher than the Cd concentration,and added ammonium ions (which reduces the pH and therefore favorsion-by-ion deposition) likewise found only the cubic phase.

Preferential (111) texture of cubic CdS on ITO/glass [14].

4.1.1.3 Mixed Hexagonal/Cubic

A mixture of phases was often reported. This is not surprising considering thesmall energy difference between them. Some examples follow.

Standard deposition giving thick films (close to 1 �m thick) on glass re-sulted in films that were ca. 90% cubic and 10% hexagonal [15].

Either not enough NH3 to dissolve the Cd(OH)2 and not stirred or a large ex-cess of NH3 and stirred [2].

Standard deposition on tin oxide/glass [16].The precipitate formed in solution was predominantly cubic, with some

hexagonal [17].


Detailed analysis of different XRD techniques led to the conclusion thatCdS used in CdS/CdTe PV cells was polytype, with essentially randomstacking of cubic and hexagonal structures in individual crystals [18].This study goes a long way to explaining the wide variation in apparentcrystal structure.

A triethanolamine/ammonia bath gave a mixed deposit [19,20].

Some comments on the role of the anion of the Cd salt are in order. Due tothe small energy difference between the two phases, small changes in adsorptionof solution species onto the growing crystals may be enough to dictate the finalcrystal structure. Early studies have shown differences in the crystal structure ofCdS precipitates, depending on the anion of the Cd salt. Halides resulted in hexag-onal CdS, sulphate gave cubic, while nitrate could give either, depending on tem-perature and pH [21].

Films deposited from a typical NH3/thiourea bath, using either the iodide orchloride salt of Cd, were studied by both XRD and ED [4]. The films deposited us-ing CdCl2 were highly textured (111) zincblende (with crystal size ca. 15 nm). Thosedeposited from CdI2 showed sharp hexagonal reflections that were not highly tex-tured. In addition, these sharp peaks rode on broad peaks, which, while not dis-cussed, suggest that most of the film is made up of a much smaller crystal size, whichmight be cubic. Using acetate or chloride anions, a well-defined peak [(probably(111)] and some much smaller peaks were found for acetate and very weak (but nar-row) peaks for chloride [22]. If ammonium salts were added (therefore lowering thepH), well-defined and strongly textured (111) peaks were obtained for both anions.

The presence of foreign ions can obviously affect the crystal structure. Morenoticeable, however, was the effect of these ions on the crystallinity. Adding Cu (byCuCl in the deposition solution) caused a decrease in the intensity of the (zincblende)XRD peaks with increasing Cu concentration [23,24]. However, there was no ap-parent change in peak width, implying that the crystal size did not change apprecia-bly, since a reduction in crystal size—whether by reduction of the coherence lengthof a fixed crystal size by defects or by actual change in crystal size—would result inpeak broadening. Similar results were obtained for films doped using CuCl2 and sub-jected to annealing in air at 300°C [25]. Doping with a variety of cations (Cu, Ag, As,In) was in all cases reported to result in loss of the XRD pattern [15].

4.1.2 Crystal Size

There are a number of factors that determine crystal size. Probably the two mostimportant are the deposition mechanism (ion-by-ion growth, in general, will resultin larger crystal size than the hydroxide mechanism, discussed in detail in Chap.3) and specific adsorption of anions onto the growing crystal (this can affect bothcrystal structure and size).

In most cases, the CdS crystal size from the standard bath was typically10–20 nm, although sometimes it could be several times larger than this, particu-


larly from lower-pH (buffered) solutions. This is larger than the typical size of theaccompanying CdS precipitate, which tends to be between 5 and 10 nm. For ex-ample, a crystal size of some tens of nanometers was deposited on a carbon-coatedTEM grid, but the precipitate in the same solution was 3–6 nm [3]. This was at-tributed to an ion-by-ion mechanism for the film vs. a hydroxide cluster mecha-nism in solution. Some large (typically 100 nm but reported up to 1 �m) hexago-nal-shaped thin crystals of Cd(OH)2 were also formed in the solution precipitate.If the growth begins by a cluster mechanism but ion-by-ion growth (whether byfree sulphide or by complex decomposition) occurs in parallel, then an intermedi-ate crystal size is a logical outcome, since the ion-by-ion growth can occur on thesmall hydroxide-formed clusters, leading to crystal growth.

The nature of the anion of the Cd salt was found to affect the crystal size insome cases, although it appears that such effects are not universal but related toother variables in the deposition process. In one report, the use of CdCl2 gave largecrystals (probably �100 nm), but with CdAc2 no XRD pattern was observed [5].From the optical spectrum of the CdS deposited from an acetate bath, a crystal sizeof ca. 5 nm can be inferred based on size quantization. Another study [22] foundfairly narrow XRD peaks (crystal size at least 20 nm) using CdAc2. CdS depositedfrom a CdI2 solution gave an XRD pattern of sharp peaks (see the previous sec-tion) on a broad background [4]. This, together with the blue-shifted optical spec-trum, suggests that most of the film is made up of a much smaller crystal size. Thecrystal size of films deposited using CdCl2 in the same study was ca. 15 nm (esti-mated from the XRD pattern).

Using a nitrilotriacetate (NTA) bath, the deposition mechanism could bemore easily controlled than from the ammonia bath; NTA is a much stronger com-plexant than ammonia, allowing pure ion-by-ion deposition if the NTA:Cd ratiois high enough. The crystal size from such an ion-by-ion deposition was �70 nm(instrument broadening limited), while from a hydroxide-mediated NTA bath itwas 5 nm [10].

A crystal size of ca. 10 nm was reported from an acidic thioacetamide bath[8]. The only other acidic bath where crystal size could be extracted was the pho-todeposition method using a thiosulphate solution, where, from the XRD, a size of�10 nm could be estimated [26]

If a comparison of crystal size and structure is made from this and the pre-vious section, a general trend appears suggesting that crystals that grow in thehexagonal modification are, in general (and there are exceptions), larger thanthose that are cubic.

4.1.3 Optical Properties

4.1.3.1 Transmission/Absorption Spectroscopy

Many studies present optical absorption or transmission spectra of the resultingfilms. (A reminder that a spectrophotometer measures transmission, not ab-


sorbance: The “absorbance” measured by a spectrophotometer is a mathematicalmanipulation of the transmission and, if reflection is present, will not be accuratewithout reflection correction. See Sec. 1.4 for details.) The purpose of these spec-tra is usually to show that the deposits are of high quality (usually interpreted tomean transparent in the subbandgap region; in most cases, scattering is undesir-able, although there may be exceptions, depending on the intended application ofthe films) and are indeed made of the material claimed (as seen from the bandgapvalue, which can be estimated from these spectra). Scattering is usually caused byoptically large (comparable to the wavelength of the light) nonhomogeneous ag-gregates; this often occurs by sedimentation of colloidal aggregates onto the sub-strate. However, it can also occur even if no colloidal phase is present in the solu-tion. There have been few studies on control of scattering in CD films. One studyreported more aggregates and lower transmission at lower deposition tempera-tures from a citrate/ammonia bath [27]. However, another study, using a tri-ethanolamine bath, reported more aggregates at higher deposition temperatures,although not in a regular manner, resulting in generally lower transmission athigher deposition temperatures [28]. Therefore, as is generally the case, such spe-cific results should not automatically be applied to all CD CdS films. Applicationof a magnetic field perpendicular to the substrate caused an increase in transmis-sion of the film [29], although it is not clear whether this is due to increased spec-ular reflectance of the field-free films or decreased scattering of the films de-posited with the field. It may be indirectly inferred from the transmission spectrathat there is no less (maybe even more) scattering in the films deposited with afield; but Atomic Force Microscope (AFM) morphology studies give the impres-sion that these films are smoother, which would be one (but not a unique) indica-tion of less scattering and greater specular reflectance.

Most CD films reported are fairly transparent to very transparent, typicallybetween 60 and 90% transparent in the subbandgap region, although lower valuesare not infrequently seen. Since the optical spectra are most often not corrected forspecular reflection, this reflection will reduce the transmission, but the “quality”of such films can be very high. In fact, a high specular reflectance is indicative of“good-quality” films, since films with a considerable degree of scattering exhibitlow specular reflectance.

The other type of information that can be extracted from optical spectra viathe bandgap is an estimation of crystal size if the semiconductor is in the size-quantized domain. This is due to the blue spectral shift caused by size quantiza-tion: The smaller the crystal size, the larger the blue shift and the larger thebandgap. This is discussed in more detail in Chapter 10. Here we note briefly somestudies where such shifts have been seen.

In most cases, the standard bath gives a crystal size larger than the largest sizethat will show an appreciable blue shift (for CdS, this value is ca. 6 nm). There aresome exceptions, however, with minor modifications of the deposition conditions.The use of iodide [4] and acetate [5] as anions of the Cd salt resulted in blue shifts


of 0.1 eV or more. Blue shifts of ca. 0.2 eV were measured in CdS films depositedfrom citrate/ammonia baths [7,27].

Transmission spectra can be modified by doping. Thus, if Cu (as CuCl) isadded to a triethanolamine/ammonia bath for CdS deposition, the effectivebandgap measured from the spectra shifted from 2.35 eV (no Cu) to just over 2.0eV [23,24]. The less steep onset and more pronounced absorption tail of theCu:CdS, together with the very low value of bandgap, suggests that this large shiftis due either to a subbandgap transition arising from Cu impurity states in the CdSor even to absorption in a separate phase of Cu-S, which would probably itself bequantized. In a somewhat similar study, but with films annealed at 300°C in air, adrop in bandgap due to Cu doping from 2.48 eV to 2.38 eV was measured [25]. Incontrast, another study reported an increase in bandgap from 2.4 to 2.48 eV upondoping with either Cu, Ag, As, or In (as well as a less steep onset). In principle,doping can affect the measured absorption spectrum in different ways. The mostobvious is introduction of levels in the gap (which would result in an apparentlowering of the bandgap). Amorphization (as seen by the loss in the XRD pattern)is commonly found to occur upon doping; the resulting disorder could cause tail-ing of the states near the band edges. This tailing would normally be seen as a de-crease in the bandgap, although an increase has been explained by splitting of thetailed levels from the bands [15]. Increase in bandgap can also occur by filling thelower-lying conduction (valence) band levels with electrons (holes), thereby re-quiring a larger photon energy to promote an electron from the valence to the con-duction band (the Burstein–Moss shift).

Antimony doping has been shown to have strong effects on the optical spec-tra [30]. The bandgap decreased from 2.47 eV (pure CdS) to 1.7 eV (nominally0.075% Sb) and then increased to 2.86 eV for 0.1% Sb. The strong drop inbandgap (the absorption was strong) for moderate Sb levels suggests an impurityband, while the increase for higher doping levels could, by itself, be explained bythe Burstein–Moss shift. However, these explanations would require effective re-moval of the impurity band at some intermediate Sb level, which would not nor-mally be expected. The crystallinity of the films does not appear (from the XRDspectra) to change appreciably with doping, thus removing amorphicity as a pos-sible explanation for the effects.

4.1.3.2 Photoluminescence

Photoluminescence (PL) is light emitted when photogenerated electrons and holesrecombine. In that sense, it is the opposite of absorption. However, while an opti-cal absorption spectrum in the great majority of cases shows the valence–conduc-tion band transition (i.e., the bandgap), photoluminescence spectra are much morecomplex as a rule. The bandgap emission (sometimes called band-to-band emis-sion) may or may not occur (it often does not), and subbandgap transitions oflonger wavelength are commonly seen. These transitions are from various surface


or bulk states in the bandgap, and they therefore can give information about thesestates, which are not seen, or at best are seen very weakly, in absorption spectra.In the majority of cases, most of the energy of photogenerated electron–hole pairsis not emitted in radiative transitions, but converted to heat in nonradiative transi-tions (i.e., the luminescence efficiency is commonly low).

It is outside the scope of this book to go into detail on the explanations forthe various PL spectra measured in CD CdS films; several examples are givenwith possible origins for the various spectral peaks.

The pH of deposition (adjusted by adding NH4OH, therefore pH increasedbut free Cd2� decreased) affected the PL spectra of the CdS films deposited froma standard solution [31]. A broad, red luminescence (ca. 1.2–2.0 eV with peak at1.68 eV) was characteristic of all the spectra, regardless of deposition pH. At pH � 11.5, a narrow (0.18 eV half-width) green peak (2.255 eV) appeared, but itdid not occur above or below this pH value). This peak, ca. 0.2 eV less than thebandgap, could be either a shallow-donor-to-shallow-acceptor transition or aband-to-fairly shallow interband state transition.

Different crystal sizes (some in the quantum size regime) were obtained byvarying the film thickness. Three interconnected PL peaks at ca. 1.83, 1.35, and 1.06eV were obtained (no green emission) [32]. A model of transitions from a deep donor(Cd–O complex) level to various other levels was suggested to explain these peaks.

Another deposition, probably from a standard solution (although the detailsof the deposition were not complete), gave a dominant peak in the green region(2.38 eV) and a broad low-energy shoulder extending to ca. 1.5 eV [33]. Decon-voluting the spectrum revealed, besides the green peak, a small yellow peak (2.25eV), attributed to a Cd interstitial–Cd vacancy (iCd–VCd) complex and red band(1.80 eV) associated with sulphur vacancies.

PL spectra of CdS deposited from two different acidic baths have been re-ported. From an acid thioacetamide bath, a broad band centered around ca. 1.5 eVwas obtained [8]. The most likely cause for this luminescence was suggested to bevalence band hole–S vacancy recombination. Films deposited under illuminationfrom a thiosulphate solution exhibited a broad band from ca. 1.46–2.0 eV (peak atca. 1.66 eV) [26].

The wide range of different PL spectra obtained shows just how much thevarious films vary from each other and the sensitivity of the (mainly surface) elec-tronic structure of the CdS to the deposition parameters.

4.1.4 Resistivity and Photoconduction of As-Deposited CdS

4.1.4.1 Dark Resistivity

Note that only as-deposited films are considered here. As will be evident to any-one familiar with semiconductor processing, annealing of these films can be car-


ried out to drastically change their electrical and photoconductive properties, de-pending on the annealing conditions. While a considerable amount of work hasbeen expended on studying the various effects of annealing on CD CdS films, thisis mostly outside the scope of this book. (See, however, the end of Sec. 4.1.7,which gives some information on the effects of rapid thermal annealing and sub-sequent removal of oxygen on the electrical properties. Also, since it is relevantfor PbS(Se) photoconductors and for photovoltaic cell use, both of which usuallyrequire some annealing, it will be treated somewhat in Chap. 5 and 9.) Two pointsare worth noting here. One is that, in general, annealing in hydrogen or vacuuminvariably reduces the dark resistivity of CD CdS, while reannealing in air or oxy-gen increases it again. Oxygen can chemisorb on the CdS surface, extracting elec-trons from the conduction band, and therefore decrease the free-electron concen-tration (hence increase in resistivity). The second point is that, for photovoltaiccell use, a lower resistivity does not necessarily mean a better cell; in fact the op-posite may even be true.

The dark resistivity of CD CdS is often, although by no means always, veryhigh. This may be reasonably attributed to the high degree of stoichiometry usu-ally obtained with CD films. This stoichiometry is certainly expected for ion-by-ion growth, and is probable also for hydroxide-mediated growth (both simple andcomplex) as long as all the hydroxide has been converted. It is likely that the caseswhere low resistivity has been reported can be explained by nonstoichiometry. Inone study [34], the activation energy of the dark conductivity (measured at andabove room temperature) was found to be �1 V, from which it can be inferred thatthe CdS Fermi level is very close to the bandgap center, meaning that the CdS ishighly intrinsic and free of common bulk defects, in particular S vacancies. Weakn-type behavior with very low donor concentration (1012 cm�3) and considerablyhigher deep trap densities (1015–1016 cm�3) were found on standard films de-posited on quartz using space charge–limited current measurements [35].

This high resistance may be responsible for the fact that the commonly usedAg contacts to the CdS behave in an ohmic manner [36], although Ag is not nor-mally considered to be a good ohmic contact to CdS in general. Two reasons canbe given for this. One is that the high resistivity of the CdS means that even an ap-preciable contact resistance may be negligible. Another factor is that, since thehigh-resistance CdS is often close to intrinsic (i.e., the Fermi level is close to thebandgap center), which by definition means a higher value of work function, evena high-work-function metal is less likely to form a Schottky (blocking) contact tothe CdS.

This is a good point to bring up briefly a property of very small crystals (of-ten obtained in CD), which is dealt with in more detail in Chapter 9. The crystalsize is, in most cases, much smaller than the size of any space charge layer thatwould be formed. This means that in an isolated nanocrystal, unless the dopinglevel is very high (usually it is not, as attested to by the high resistivities more of-


ten obtained), there will not be much built-in electric field. The situation of a filmof aggregated nanocrystals is not so obvious, but it is likely that a space charge layerof the normal type will not be obtained. This has important implications in consid-ering contacts both between nanocrystal and metal contact and between nanocrys-tals. As far as (photo)conducting properties are concerned, terms such as Schottkycontact or grain boundary barriers need to be considered with this point in mind.

Values for the resistivities of various CD CdS films are given in Table 4.1.The first thing that can be observed is the lack of any obvious correlation betweenthe resistivity and deposition conditions in most cases. Only some conditions aregiven in the table; the details of the concentrations are not always available; evenif they are, it would be oversimplistic to try to compare them based only on con-centration. For example, the ratio of Cd salt to complexant is no less (probably

TABLE 4.1 Dark Conductivity and Photoconductivity of CD CdS Films

Resistivity (�-cm)

Bath conditions Dark Light Reference

CdCl2 90�C 109 700 38CdSO4 70�C 108 1000 39CdAc2 or CdCl2 80�C 106–107 103–104 5CdAc2 104–106 10–1000 16Cd(NO3)2 room temperature 1012 Small effect 40CdSO4 90�C 104 — 11(80–95�C) pH 10–11 340–600 — 22(80–95�C) pH 9–10 15–150 22CdCl2 106–108 102–104 4CdI2 108–1010 103–105 4CdAc2 NH4

� 85�C 103–104 — 41Triethanolamine 26�C 109 1000 28, 36Triethanolamine 75�C 109 1 28, 36Citrate/ammonia 60�C 5.107 2000 27Citrate/ammonia 75�C 108 5 27Citrate/ammonia 50�C 2 � 108 200 7Citrate/ammonia 70�C 2 � 108 4 7Thioacetamide (pH � 8) 40�C 104 — 37

Temp. of minimum resistance 63�C 10 —Temp. of minimum resistance 78�C 400 —

CdCl2 85�C No magnetic field 2 � 105 — 42CdCl2 85�C Magnetic field 60 —

The first column gives the Cd salt and deposition temperature. If not specifically noted, the bath is astandard ammonia/thiourea bath.


even more) important than the absolute concentrations themselves. Also, so manyparameters are interrelated.

In an attempt to say something intelligent about these resistivities, there ap-pears to be some correlation between the pH and resistivity, with low resistivityobtained when the pH is relatively low (only a few experiments have been carriedout at relatively low values of pH; also note Ref. 22, which describes an anoma-lously low resistivity even at “normal” values of pH). The bath described by Itoand Shiraishi [37] is very different from the previous ones, for three reasons: therelatively low pH (� 8), the use of thioacetamide instead of thiourea, and the flowsystem used in this deposition. Very low values of dark resistivity were obtainedwith this bath and with an unusual temperature dependence (a minimum of 10 �-cm was found at 63°C, which increased on either side of this temperature value).It was suggested that Cl, from the NH4Cl buffer, acted as a dopant; however, otherchloride baths gave much higher resistivities.

Some weak correlation between film morphology and resistivity was notedfor films, deposited from a pH 9.5 bath: Films deposited from a closed system (noevaporation or loss of ammonia) were more uniform and had somewhat lower(two to three times) resistivity (ca. 1.5 � 104 �-cm) than films deposited from anopen bath [43].

The triethanolamine bath gives consistently high resistivities, independentof the deposition conditions, and the citrate/ammonia bath behaves similarly, al-though with somewhat lower resistivities. It should be noted that chemisorption ofoxygen on the CdS is known to increase the resistivity, and some (many?) differ-ences may well be due to different surface chemistries of the CdS crystals. Thus,by definition the complexant can bind to surface Cd and cover the surface in somecases. This can lead to (at least) two opposing effects: “insulation” of one crystalfrom another and prevention of oxygen chemisorption (this latter need not, how-ever, necessarily lead to lower resistivity but the opposite if the adsorbed complexacts in a similar way to the oxygen). This surface adsorption may explain, e.g.,both the high resistivity of the triethanolamine-bath films and their relative inde-pendence from other deposition conditions. It is worth noting that most (althoughnot all) CD CdS films have a rather similar crystal size; therefore this factor doesnot seem to be as important as might have been expected.

Films deposited from a standard bath but under application of a magneticfield perpendicular to the substrate exhibited resistivities that decreased stronglyas the strength of the magnetic field increased [42]. The resistivity, which for afilm deposited without the field was 2 � 105 �-cm, was as low as 60 �-cm on ap-plication of a 77-mT magnetic field. From the dark resistivity/temperature values,a level with an activation energy of 0.68 eV was found for the field-free films anda shallow level of 0.053 eV and others at �0.17 eV for those deposited in the pres-ence of the field. The shallow level at 0.053 eV was suggested to be due to excessCd. It was suggested that the magnetic field might affect the rate of arrival of cad-


mium and sulphide ions at the substrate and that more sulphur vacancies areformed in the presence of the field, hence the lower resistivity (see Sec. 4.1.6.8 formore details on this deposition).

4.1.4.2 Photoconductivity

CD CdS films are usually strongly photoconducting as deposited. This is in con-trast to most other CD films, which tend to be only weakly or moderately photo-conductive in the as-deposited state [44]. Values of the resistivity under illumina-tion are given in Table 4.1 where available. Most experiments have usedillumination intensities comparable to solar irradiation.

As is the case for the dark resistivity, the dependence of the sensitivity of thephotoconductivity (defined here as the ratio between light and dark conductivity)on the deposition parameters is far from clear-cut. Some observations can bemade, however. The first (obvious) one is that for a high sensitivity, the dark re-sistivity must be high. Apart from this, there does seem to be a general trend(clear-cut in the triethanolamine and citrate baths and seen also by the lack of ap-preciable photoconductivity in the one low- (room-) temperature-deposited filmreported [40]) of an increase in photosensitivity (due to decrease in light resistiv-ity) with increasing deposition temperature.

For the standard baths, the sensitivity varies in most cases between 103 and106. A CdI2 bath resulted in somewhat greater sensitivity (as well as dark resis-tivity) than a CdCl2 bath [4]. The deposition temperatures of these two baths weredifferent, but it was reported that the film properties were independent of the de-position temperature.

Another study found no appreciable difference in either dark or light resis-tivity between acetate and chloride baths [5]. Interestingly, there was apparently alarge difference in crystal size between the two baths (see Sec. 4.1.2 on crystalsize), which implies that the crystal size is not an important factor in determiningthe resistivity or photoconductivity, at least for this bath.

The triethanolamine bath showed a distinct trend of the photosensitivitywith deposition temperature [28,34,36,45,46]. The photosensitivity was higher(which, since the dark resistivity was temperature independent, means the light re-sistivity was lower) for higher temperatures, with the major change occurring be-tween 30 and 45°C.

The citrate/ammonia bath has much in common with the triethanolamine/ammonia bath with high light sensitivities, particularly at higher deposition tem-perature, and little temperature dependence of dark resistivity. This is in spite ofthe very different bath compositions and concentrations. In particular, the citratebath contained much lower concentrations of Cd and, as a result, was more highlycomplexed.

One study showed a very strong dependence of dark resistivity on measur-ing temperature and a much weaker dependence of the resistivity under illumina-


tion, with the result that the sensitivity was very temperature dependent [15] (seeSec. 4.1.4.3).

The decay time of the photoconductivity, �, is another important parameterfor which, as with the dark resistivity and photoconductivity, there is often no ob-vious correllation with the deposition parameters. In many cases, the decay is veryslow (hours), particularly for the triethanolamine bath, where it can be greater than10 hr. Decay times of hours have been reported for the standard bath [4,16], al-though if CdI2 was used instead of CdCl2 in the latter, � dropped to seconds (thedeposition temperature was different for the chloride and iodide bath, but report-edly this did not affect the properties). For films deposited from the citrate bath,temperature-dependent decay times from �1 min (deposition at 75°C) to tens ofminutes (60°C) were reported in one study [27] and tens of minutes with only asmall deposition temperature dependence in the other [7]. Differences in the de-position conditions of these two studies were described earlier. From the tri-ethanolamne bath, much longer decay times were observed at higher depositiontemperatures, as shown in Figure 4.1 [34]. Another study, using ammonium-buffered standard films (possibly ion-by-ion deposition), found decay times ofseconds for single films, which increased to several minutes for multiply de-posited films [47]. The photosensitivity of these latter films was less than for mostothers (ca. 30 for the single films and �10 for the multiple ones).

FIG. 4.1 Time dependence of photoconductivity of CdS films deposited from tri-ethanolamine/ammonia/thiourea bath at two different temperatures (26°C and 75°C). Thetwo plots at each temperature differ by the ratio between the Cd and thiourea concentra-tions: [Thiourea]:[Cd] � 0.25 for the upper plots at each temperature and 0.5 for the lowerplots. (Adapted from Ref. 34 with permission from Elsevier Science).


It is important to note that oxygen adsorption has a strong effect on the pho-toconductivity decay of the triethanolamine films [36]; it is probable that this isnot limited only to these films. Oxygen greatly decreases the decay lifetime, asseen by the increase in � for aged films compared with freshly deposited ones. Inthis same study, the photoconductivity sensitivity and decay time both decreasedgreatly with increase in measuring temperature (the former due mainly to the ac-tivated decrease in dark resistivity with increase in temperature, but also duepartly to decrease in resistance under illumination). The possible effects of oxy-gen on photoconducting parameters are discussed in this reference (general back-ground on photoconductivity is given in Ref. 34). Here it is enough to note that ad-sorbed oxygen is believed to extract electrons from the CdS conduction band andalso to introduce deep trapping centers (interband surface states) that increase thecarrier lifetimes, thereby increasing the photoconductivity decay time. Oxygencan also affect electron mobility between grains by modification of grain bound-ary barriers between crystals.

A comparison was made between films deposited from standard baths usingeither CdCl2 or CdAc2 [48]. While for most measurements these films were an-nealed at 300°C, and therefore are not compared with the as-deposited films here,thermally stimulated current (TSC) measurements were carried out on as-deposited films. Such TSC provides an indication of the density and energy oftrapping centers: The magnitude of the current, obtained by heating the sampleand exciting charges out of traps into a band, is an indication of the trap densityand the temperature at which the current is generated is a measure of the trap en-ergy. The trap density was much higher and the traps considerably deeper for theacetate-prepared films than for the chloride ones (after annealing, the trap densitywas higher for the chloride films, seen as a large increase in TSC for the chloridefilm after heat treatment).

We can conclude this section with the insight, gained from this overview ofthe electrical and photoconductivity properties of these films, that, in spite of themany studies already carreid out, a comprehensive and systematic study of theseproperties and their correlation with a wide range of deposition parameters is stillneeded in order to understand what determines these properties. These studiesshould also include postdeposition treatments—not so much annealing, which hasbeen carried out, but surface treatments (e.g., immersion in triethanolamine),which could show the importance (or lack of it) of the crystal surface condition.

4.1.4.3 Electrical Properties of Doped CdS Films

Doping can be divided into two parts: native doping (e.g., S vacancies) and ex-trinsic doping by foreign elements. This section deals with the latter, not becauseit is more important but because there is little in the literature to link native dop-ing with the electrical properties of CD films. It will be enough to note that the fewmeasurements of ND (donor density) carried out tend to give values typically


around 1016 cm�3 (e.g., Ref. 49 for epitaxial, hexagonal CdS on (1̄ 1̄ 1̄) InP fromC–V measurements) or even larger. This is somewhat surprising, since, if thesefilms are so highly stoichiometric as expected, lower values might be expected.Two comments here. One is that the doping may originate from the surface, sincethe surface-to-volume ratio of the CdS crystallites is high. The other, also relatedto the high surface area, is that errors in interpreting C–V measurements com-monly used to derive ND may arise because of the lack of knowledge of the truesurface area.

We will now consider individual dopants used in CD CdS.

Boron. Boron, substituted for Cd, is a donor in CdS. B-doped films weredeposited by adding boric acid to a standard deposition bath, with the B:Cd ratiovarying from 10�5 to 10�2 [16]. The boron was assumed to occur in the form ofborate ions (BO3

2�). The dark resistivity dropped nearly three orders of magnitudewith optimum B content (B:Cd ratio in solution of 0.001), from 2.104 to 30 �-cm.At higher B concentrations, the dark resistivity again increased until, at a B:Cd ra-tio of 0.01, the original resistivity of the undoped CdS was regained, and it did notchange with increased B content. However, the resistivity under illumination de-creased to ca. 3 �-cm, almost independent of the B content. Thus high B contentincreased the photosensitivity of the CdS, although only by a factor of 2–3. Therate of decay of the photocurrent was greatly reduced by B doping, from ca. onehour for undoped films to as much as several tens of hours for doped ones. Thissuggests a deep trap resulting from the B, separate from the shallow donor that isresponsible for the drop in resistivity.

Nitrogen. Nitrogen ions (N� with energy of 130 KeV) were implantedinto CdS deposited from a triethanolamine/ammonia bath [20]. The resistivity ofthe as-deposited films was ca. 108 �-cm and dropped, depending on ion dose, upto seven orders of magnitude for an ion dose of 1017 ion/cm2. Even more notable,the conductivity type changed from n-type (the normal type for CdS) to p-type, asmeasured by hot probe. An acceptor level, 0.6 eV above the valence band, was in-troduced by the ion implantation.

Copper. Copper was doped into triethanolamine/ammonia films byadding CuCl in the deposition solution [23,24]. Resistivity dropped from 109 �-cm (undoped) to a minimum of 0.5 �-cm for optimum doping (the Cu content ofthe CdS was not measured) and the conductivity was p-type. Both higher andlower Cu concentrations in solution gave higher resistivities; it was surmised thattoo high concentrations of Cu in the deposition solution resulted in rapid precipi-tation of the Cu as CuxS, depleting the solution of Cu. While the high-conductiv-ity p-type CdS was, not surprisingly, not photoconductive, films with smalleramounts of Cu were photoconductive (although the photoconductivity gain wasless than for nondoped CdS, and the response time of the photoconduction, both


rise and fall, was much shorter than for the undoped film). The Cu-induced statesresulted in recombination centers rather than long-lived trap states characteristicof CD CdS films. Another study of Cu-doped CdS films reported different results,although there were differences in the preparation; in particular the films were an-nealed in air at 300°C, and CuCl2 (rather than CuCl) was used [25]. In this work,dark resistivity did not vary greatly for low Cu concentrations (ca. 107 �-cm), butincreased by nearly two orders of magnitude for high Cu concentration. Also, thephotoconductivity response increased with Cu content. The annealing carried outin this study calls for caution in any comparison with the previous one. A thirdstudy (no variation in Cu content) found a decrease in dark resistance (from ca.108 to ca. 104 �-cm upon Cu (also Ag, As, and In) doping [15]. The dark resis-tivity was very highly temperature dependent, especially for the undoped CdS (ca.106 �-cm at 35°C and 104 �-cm at 50°C). Since the resistance under illumination(�105 �-cm for undoped and �104 �-cm for doped) was much less temperaturesensitive, the photoconductivity response was very temperature dependent, beingmore pronounced at lower temperatures.

Lithium. Lithium can act as an acceptor in CdS. Shikalgar and Pawar stud-ied electrical [50] and photoconducting [51] properties of Li-doped CdS [standarddeposition with 0.1% (by weight of CdS) Li salt added to the deposition solution].The addition of Li increased the dark resistivity by a factor of 3–4 (resistivitieswere given as ohms and the exact geometry of the measuring system was not de-scribed: however, specific resistivities could be estimated to be of the order of 107

�-cm). The room-temperature energy of activation for both doped and undopedfilms was ca. 22 meV, i.e., shallow donor conduction. Above ca. 60°C, the resis-tivity dropped much more rapidly as a function of temperature, with an activationenergy of ca. 1.2 eV (intrinsic conductivity). In addition, the Li-doped film ex-hibited an intermediate level at 0.16 eV in the temperature range of ca. 40–65°C,ascribed to a Li acceptor level. The Li-doped films, like the undoped ones, weren-type; CdS is difficult to dope p-type due to self-compensation, and since the re-sistivity of CD CdS films are normally very high, it is not surprising that acceptordoping does not increase this resistivity very greatly.

The light:dark conductivity ratio of these Li-doped films was not explicitlygiven, although an order of magnitude value of 106 could be inferred from the re-sults. The photocurrent–time behavior for the CdS:Li (the equivalent data for theundoped films were not given) was history dependent. Initially, the photocurrentincreased linearly with time (over a maximum measured time of 10 min), but in-creased more rapidly and exponentially with larger photocurrents after light–darkcycling. In all cases, the decay was multiexponential and slow, typically tens ofminutes. These measurements were carried out in vacuum; if air was introduced,the steady-state photocurrent decreased, attributed to oxygen adsorption on thesurface of the CdS crystals, resulting in extraction of electrons from the conduc-


tion band. Interestingly, the peak of the photocurrent of the CdS:Li was blue-shifted (480 nm) compared to the undoped CdS (520 nm); no explanation couldbe given for this effect.

Aluminum. Aluminum, as a trivalent ion, should be an n-type dopant forCdS. A small decrease in resistivity (by a factor of 2) to ca. 103 �-cm was foundwhen Al2(SO4)3 was added to a standard bath [52]. In another study, Al was addedas Al2(SO4)3 to a thiourea bath (85°C) of relatively low pH (9.5) [53]. The resis-tivity of the undoped CdS was ca. 3 � 105 �-cm and decreased at least an orderof magnitude on doping with Al. Codoping with chlorine (as NH4Cl in the solu-tion) decreased the resistivity almost another order of magnitude. In both studies,an excess of Al resulted in an increase of resistivity. This was explained by excessAl3� occupying interstitial positions: however, interstitial Al3� might be expectedto increase the n-type doping and therefore decrease the resistivity, and it is notclear why the resistivity should increase. It may be that an insulating Al(OH)3

phase occurs if too much Al is added.

4.1.5 Substrate-Dependent Growth and Epitaxy

4.1.5.1 Introduction

There are a number of studies that report the effects of the substrate on the CdSfilms. With the exception of epitaxial deposition, which will constitute the mainpart of this section, it is usually difficult to explain any specific substrate effect.Also, it should be borne in mind that each specific study is confined to one depo-sition bath and that a substrate effect obtained for one bath need not necessarily beobtained for a different one.

Some examples of substrate effects on the film can be given. Strong (0002)or (111) texturing was obtained on glass substrates but much weaker texturing onSnO2/glass [16]. Much poorer crystallinity (this may also mean smaller crystalsize) of the CdS was obtained on Si than on glass or ITO/glass [54]. Using XRDpeak shifts and optical absorption spectroscopy, the presence of strain in as-grownCdS on both glass and ITO/glass was inferred [55]. The strain was greater for thefilms deposited on the ITO, and this was attributed to mismatch strain between theCdS and ITO.

4.1.5.2 Epitaxy

Various investigations into the epitaxial deposition of CdS onto different single-crystal substrates have been carried out by Lincot et al. On InP, which is closelylattice matched to CdS (�0.1% difference), epitaxial deposition (c-axis of hexag-onal CdS perpendicular to the substrate) occurs on the (1̄1̄1̄) P polar face of the InPbut polycrystalline deposition on the (111) In face [49,56]. This difference wasclearly due to differing chemical or electrostatic interaction between the InP faces


and the constituents of the deposition solution, since the lattice spacing is the samefor both polar faces. The degree of epitaxy on the P face was also dependent onthe deposition conditions: in particular, higher temperatures resulted in better epi-taxy, as might be expected due to the higher mobility of the depositing species onthe surface. The epitaxy was maintained up to at least 100 nm film thickness (themaximum thickness studied).

On (100) etched (Br-MeOH) InP, cubic, fairly well-oriented CdS wasformed, although with many small-angle grain boundaries. In the absence of theetch treatment (using only H2SO4 to remove native oxide, as was also employedfollowing the Br-MeOH treatment), only polycrystalline CdS was deposited,which showed that not only the crystal face, but also the manner in which that facewas pretreated, is important [57]. The formation of the cubic phase in contrast tothe hexagonal phase formed on(1̄1̄1̄) InP was attributed to the lack of lattice matchbetween the (cubic) (100) face of InP and hexagonal CdS.

GaP has a much larger mismatch with CdS (�7%) compared with InP. Yeta fair degree of epitaxy was obtained for CD CdS on the (1̄1̄1̄) GaP surface [58].In this case, a mixture of cubic and hexagonal CdS with a large density of stack-ing faults, presumed due to strain relaxation arising from the large mismatch, wasobtained.

Because of the importance of the junction between CD CdS and CuInSe2

(CIS) for thin-film photovoltaic cells (see Chap. 9), as well as the relatively smallmismatch between CdS and CIS (�0.7%), deposition onto oriented CIS films hasalso been studied [59,60]. Two different CIS faces were studied—(100) and (112).As with the (100) P face of InP, because there was no lattice match between thisface and any hexagonal CdS face, cubic, epitaxial CdS was deposited. On (112)CIS, which matches either (111) cubic or (0001) hexagonal CdS, a mixture of bothphases was deposited, with a moderate degree of epitaxy, which improved if theCIS was first subjected to a cyanide treatment (cyanide removes excess CuxSe andvarious Se species and generally cleans up the surface). The epitaxy also improvedwith increase in temperature. The transition temperature was quite abrupt: Below60°C, the films were polycrystalline, while above this temperature they were epi-taxial, with increasing perfection as the temperature increased. Since the deposi-tion at lower temperatures was much slower than at higher ones (for a temperaturedifference of ca. 40°C, the deposition rate increased 30 times), this implied thatthe increased mobility of the depositing species on the surface was not necessar-ily the main factor in the temperature dependence of epitaxy. It was suggested thatfaster and more complete decomposition of reaction intermediates was an impor-tant factor in determining the epitaxy. If so, this is presumably true for epitaxialdeposition on other substrates.

The composition of the deposition solution was important in order to obtainepitaxy on CIS, which, in contrast to the epitaxy on InP, was not obtained using“standard” solutions. Instead, low Cd concentrations (maybe more important, low


Cd:NH3 ratios) and lower than usual pH (adjusted by adding ammonium ions)were necessary. These are the factors that favor an ion-by-ion mechanism overone involving Cd(OH)2, which may explain the need for this modified solution(most commonly used solutions operate in or near the region where Cd(OH)2 canexist, at least on the substrate). This explanation still leaves open the question ofwhy good epitaxial growth can be obtained on InP under some conditions but notCIS from a “standard” solution. While there is no present answer to this question,considering the sensitivity of the epitaxy to the chemical properties of the surface(such as etch or differences between P and In polar faces of InP), this should notbe too surprising.

The epitaxy was maintained for CdS thicknesses up to 100 nm, after whichthe deposit became polycrystalline. This transition coincided with the visual for-mation of CdS in the solution, which resulted in a switch of the mechanism froman ion-by-ion growth, necessary to obtain epitaxy, to one involving colloidalspecies. Since, in principle, conditions can be chosen so that only an ion-by-iongrowth occurs, it can be expected that much thicker epitaxial films are obtainablefrom CD on suitable substrates.

4.1.5.3 Deposition on Monolayers: Selective Growth and Patterning

If a substrate is not clean, films either do not grow or grow with poor adherenceon the “dirty” parts of the substrate. This has been exploited by partially coveringthe substrate with a monolayer. When a mica substrate was incompletely coveredby a monolayer of octadecylphosphonic acid, CdS growth was found to occurpreferentially on the mica [9]. This was shown also for CdS deposited on an oc-tadecyltrichlorosilane- (OTS)-coated Si substrate and was used to pattern the CdSdeposit by applying the OTS onto the Si using a patterned stamp [61]. Either theOTS could be removed by sonication or, even without removal of the OTS, depo-sition occurred only on the bare Si if the CdS was not too thick (ca. 50 nm). Edgeswith a variation of ca. 100 nm could be deposited by this method.

4.1.6 Variations in Preparation

4.1.6.1 Variation in pH

While deposition rate normally increases with increase in pH for the standardbath, using an ammonium salt to lower the solution pH resulted in the opposite be-havior; i.e., increased pH led to slower deposition [22]. The pH in these experi-ments was increased by adding NH4OH. Increased pH (and NH4OH) results intwo opposing effects: Thiourea decomposition increases but free [Cd2�] de-creases. Since the deposition rate for the solution with no added NH4

� increaseswith increase in pH, the former apparently outweighs the latter. When extra am-monium ion is added, much more ammonium hydroxide is needed to increase the


pH to the original value, and the decrease in free [Cd2�] dominates the reactionkinetics. More generally, addition of an ammonium salt (apart from the hydrox-ide, which increases the pH) increases the complexation of Cd without increasingthe pH (it actually decreases the pH of an NH4OH solution due to buffer action).This results in a decrease in deposition rate due to a lower free-Cd2� concentra-tion and, if the pH decreases, reduction in the thiourea decomposition rate. Aslower deposition allows the formation of thicker films, since less CdS will beformed homogeneously in the solution. This has been shown in many studies (e.g.,Refs. 22, 40, and 43). In addition to slower thiourea decomposition, lower pH willalso decrease the likelihood of Cd(OH)2 formation and will therefore favor (rela-tively) an ion-by-ion rather than cluster mechanism.

Only a few acidic baths have been described (see also sections 4.1.6.6 and4.1.6.7 below). In one, thiosulphate was used at a pH of between 2 and 4 and a tem-perature of 85°C [131]. The bandgap was 2.35 eV and the resistivity 104–105 �-cm.Thioacetamine has also been used at a pH of 5 [8]. The films from this bath wereclearly hexagonal. The rationale for using an acid bath is to prevent the formationof hydroxy species; this is a major problem for ZnS but much less so for CdS.

4.1.6.2 Variation in Complexant

Cyanide, a stronger complexant than others used, has been employed as a com-plex for CdS deposition [62]. Except for the fact that thicker films could be ob-tained (ca. 1 �m compared to a few hundred nanometers with the standardmethod), the properties of the films made with cyanide [crystal structure, crystalsize, bandgap (measured to be 2.2 eV, an anomalously low value but the same asthat of films deposited from a standard bath in the same study)] were the same asthose of the standard bath. Solution composition details were not given.

Films up to 3 �m thick were obtained with the triethanolamine/ammoniabath [19]. It is probable that this larger-than-normal thickness is due to depositionoccurring via an ion-by-ion mechanism, due to the additional complexing by thetriethanolamine, the somewhat lower pH than usual (10), and, for the 3-�m film,the low deposition temperature employed (30°C), factors that reduce free-Cd2�

and/or OH� concentrations, thereby favoring the ion-by-ion mechanism. Thiswould result in thicker films since no (or at least less) CdS is wasted as a homo-geneous precipitate in the solution.

Ethylenediamine has been used as a complexant [6]. It is a stronger com-plexant than ammonia and therefore only needs to be used in low concentration(between two and four times that of Cd).

4.1.6.3 Variation in Thiourea Concentration

Increase in the concentration of thiourea clearly leads to an increase in depositionrate. Additionally, it has been seen that the defect density (measured from TEMmicrographs as structural defects such as stacking faults) decreased greatly with


an increase in thiourea concentration from 1011–1012cm�2 for a thiourea concen-tration of 28 mM to 1010cm�2 for a concentration of 100 mM [12].

4.1.6.4 Anion Effects

Some effects of anions have been noted previously. These are often not consistentand, in general, it is difficult to attribute the effect of the anion of the Cd salt to anyspecific effect with any degree of confidence. Two studies on the effect of the anionon the rate of deposition did find small but significant differences (a factor of 2),which more or less were in agreement [63,64]. The latter found the rate to increasein the order: CdI2, CdSO4, Cd(NO3)2, Cd(CH3COO)2, and CdCl2; this series corre-sponds approximately with the decrease in (negative) electrode potential and corre-spondingly to the decreased strength of complexation between Cd2� and the re-spective anion. Additionally, the rate decreased with increase in concentration of theanion (added as an alkali metal salt). These observations suggest that the effect maybe due to mild additional complexation compared to that of ammonia alone.

4.1.6.5 Surfactants

The addition of surfactants to the standard CdS bath resulted in a reduction in therate of deposition and an increase in the terminal thickness [65]. Surfactants adsorbonto surfaces (both the substrate and colloidal particles in the solution), and there-fore it is not surprising that the growth rate is reduced. The adsorption of the sur-factant onto CdS colloidal particles also can prevent flocculation and precipitation,thereby increasing the CdS available for deposition (hence, presumably, the in-creased terminal thickness). At the same time, it is possible that the presence of sur-factants (or any strongly adsorbed species) might prevent sticking of the colloidalparticles to the substrate and to each other, in the same way as they prevent floc-culation, which is exactly the sticking together of the colloids. However, there isno evidence that this actually occurs in this study. Another effect of the surfactant,inferred from the slightly higher than usual bandgap (2.52 eV), is the small size ofthe CdS crystals; although not measured, it can be assumed to be ca. 5 nm from thebandgap shift, presumably due to size quantization (see Chap. 10). This is not sur-prising since crystal growth is in competition with adsorption of the surfactant.

4.1.6.6 Electrochemical/Chemical Deposition

Yamaguchi et al. described an interesting extension of the CD process for CdS us-ing a parallel electrochemical step [66]. They termed this process electrochemi-cally induced chemical deposition. It is based on electroreduction of protons in so-lution, which results in an increase in pH locally at the electrode. They usedthioacetamide as a sulphur source. In the acid solutions in which the deposition iscarried out (pH between 1.6 and 4.6), no film deposition of CdS occurs (althoughit does precipitate in the solution) in the absence of the electrochemical proton re-duction. In the presence of proton reduction, CdS films were formed. These films


were characterized by relatively large, hexagonally faceted wurtzite CdS crystals(crystal sizes from a few tens of nanometers up to 300 nm, the larger sizes beingformed at lower pH). The CdS precipitated in solution was quite different; it wasnot faceted and had smaller crystal size (ca. 15 nm). It was suggested that the filmgrowth proceeded by a surface-catalyzed decomposition of a Cd–thioacetamidecomplex and that the electrochemical proton reduction affected the surface prop-erties (presumably the surface of both substrate and growing CdS) in such a wayas to reduce the activation energy needed for the deposition reaction. The growthwas an atom-by-atom (or ion-by-ion) process, leading to larger crystal size thannormally obtained by the hydroxide-mediated particle growth.

By adding a strongly adsorbing species (2-mercaptoethanol) to such a de-position bath, they were able to reduce the crystal size by varying amounts due tosurface capping of the growing crystals, preventing further crystal growth but al-lowing nucleation to proceed. Thus, the film thickness was not strongly affectedprovided the mercaptoethanol concentration was low (�10 mM); above this con-centration, film growth was prevented, as would be expected, since adhesion be-tween crystals and substrate or between different crystals would probably be poorin the presence of an adsorbed coating [67]. The resulting nanocrystalline filmsexhibited quantum size effects (see Chap. 10 for more details).

4.1.6.7 Illumination-Induced Growth

There have been a few experiments related to the effect of illumination of thegrowth of CdS films. Simple heating of the deposition bath by absorption of theradiation is one obvious factor that can affect the deposition [68]. However, evenin this case, other effects occur, since the color of the bath was reported to darkenif UV (sunlight) illumination was employed. Based on previous studies of illumi-nated CdS colloids when elemental Cd was formed, both as a film and in solution[69], as well as the known tendency of ZnS to undergo reduction to metallic Znunder UV illumination, this darkening may be assumed to be caused by elemen-tal Cd. There are several possible mechanisms that may explain such an effect; re-duction of the CdS by photogenerated electrons is one possibility.

A variant of CD was based on illumination of a solution containing thiosul-phate and cadmium ions by UV light [26,70,71]. CdS was deposited only on theilluminated portion of the substrate. Since only light absorbed by thiosulphate(wavelength shorter than 300 nm) was effective, the effect was attributed to pho-todecomposition of thiosulphate to elemental S and solvated electrons and subse-quent reaction with Cd2�.

S2O32� � h�D S � SO3

2� (4.1)

2S2O32� � h� → S4O6

2� � 2e� (4.2)

SO32� � S2O3

2� � h� → S3O62� � 2e� (4.3)

Cd2� � S � 2e� → CdS (4.4)


The optimal pH was ca. 3.5 (at lower values, the film was contaminated with ele-mental S, which forms spontaneously in the dark), while the deposition rateslowed down at higher values). The CdS was sphalerite, and, from an examinationof the XRD, a crystal size of �10 nm could be estimated. The bandgap was 2.42eV, the literature value for CdS.

4.1.6.8 Deposition Under the Influence of an ExternalMagnetic Field

Deposition has been carried out from a standard bath with a magnetic field ap-plied, both parallel and perpendicular to the substrate [29]. Differences werefound in the film properties only for a field perpendicular to the substrate. Thetransmission of the films in the nonabsorbing region was ca. 10% higher (see Sec.4.1.3.1). The films deposited with the field were up to three times thicker thanthose deposited using the same conditions but in the absence of a field. The darkresistivity of the field-applied films was much less than that of the field-free ones(see Sec. 4.1.4.1). The cause of these effects is not clear.

Similar measurements were carried out using an external electric field [72].Some differences in morphology and optical properties were measured, depend-ing on the direction of the field with respect to the substrate. It is not clear, fromthe experimental setup, why the field should influence the deposition, since thefield is external and should drop across the air and the glass walls of the reactionvessel.

4.1.6.9 Deposition on/in Porous Silicon

Porous silicon is under extensive study, largely due to its luminescence properties.For electroluminescence, however, some form of contact has to be made with theSi, and this necessitates deposition of another phase inside the pores of the Si inorder to contact as much as possible of the internal area of the high-surface-areaSi. With this in mind, CdS has been deposited inside the pores of porous siliconvia a two-stage method [73]. Cd(OH)2 was deposited from an ammoniacal bath atpH 8, followed by conversion of the Cd(OH)2 to CdS by treatment with thioac-etamide at pH 8. This was repeated several times until the pores were essentiallyfilled with CdS. The reason that this two-stage process was needed is that eitherthe Si was unstable at the temperatures and pH values needed to deposit CdS froma thiourea solution, or CdS was formed in solution rather than on the Si surface us-ing thioacetamide.

4.1.6.10 Bath Geometry

One of the disadvantages of the CD process as usually carried out is the largewaste of materials (for example, in CdS deposition, most of the Cd—often over90%—is unused in the film deposition because it deposits homogeneously in so-lution and/or on the walls of the reaction vessel). Probably more important than


the material loss is the environmental concern of disposing of this Cd (or otherheavy metal) if it is not recycled. The same goes for the ammonia often used as acomplexant—much is lost to the atmosphere and, in common with other com-plexants, much is wasted to tie up the heavy metal ions.

Use of low concentrations of metal ion (ca. 1 mM) presents a partial solu-tion to this problem. However, for any industrial process, a continuous-flow sys-tem seems the best option. Ito and Shiraishi flowed a solution of thioacetamideand a Cd salt into a 0.5-mm-thick flow space [37]. A detailed flow system hasbeen described for CdS by Boyle et al. [74]. There are several features in this sys-tem, shown schematically in Figure 4.2. Probably the most important is the locallyheated substrate. Since CD reactions are usually very temperature dependent, byheating only the substrate (in this case, by resistive heating), deposition is limited,to a large extent, to the substrate. This system also uses ethylenediamine insteadof ammonia, which greatly decreases loss by evaporation as occurs with ammoniain an open system. Filters are employed in the flow system to remove any colloidalmatter formed. Fresh reagents can be added, as required, to the recirculatingclosed-loop system. In connection with the foregoing studies, a batch process,whereby the solution was filtered after deposition and complete reaction andreused (with the addition of more reactants as required), was also shown to be fea-sible [75]. While the deposition rate slowed down for successive depositions, thiscould be compensated for by increasing the concentrations of various reactantsfrom run to run. The photovoltaic parameters of Cu(In,Ga)Se2/CdS solar cells fab-ricated using this approach was not found to vary from deposition to deposition.

FIG. 4.2 Flow system for continuous deposition using a locally heated substrate. (AfterRef. 74 with permission from Elsevier Science.)


By minimizing the spacing between substrates, reagent utilization can bemaximized by increasing the ratio of substrate surface to solution volume [27].The maximum film thickness was reduced if the substrate spacing was too close—depending on the deposition parameters, and this was related to a “critical thick-ness” of reagent layer at the substrate surface connected with the presence of col-loidal particles in the solution. Typically some 1- to 10-mm spacing was necessaryto obtain the maximum thickness (200–400 nm in this case). This study was car-ried out under conditions where a hydroxide cluster mechanism was operative (thesolution was already turbid in the early stages of deposition), and the results can-not be extrapolated to other mechanisms.

4.1.7 Impurities in Chemical Deposition CdSThere have been a number of studies involving impurities in the CdS films, withvarious results. It must be emphasized that if the films are not very well rinsed af-ter preparation (and possibly even if they are), some of the ions involved in thepreparation may be present as adsorbed species. The most comprehensive study,involving a range of analytical techniques, sums up the probable impurity situa-tion for films deposited from standard baths [76].

The main impurity, not unexpectedly, is oxygen (ca. 11 atomic %). Evi-dence was presented to show that this O was probably mainly in the forms of car-bonate and adsorbed water. The carbonate could come from two sources: dissolu-tion of atmospheric CO2 and (see Eq. (3.11)) from decomposition of thiourea.

Nitrogen (ca. 5 at. %) occurs as carbon–nitrogen bonds, probably mainlycyanamide (NCN2�), although other C–N bonded compounds were also believedto be present. If cyanamide is present as the Cd salt, this would tie up 5% of theCd. The Cd:S ratio was found to be only slightly higher than unity (ca. 1.02), andsome of the Cd may be bound to carbonate. Therefore other C–N species are likelyalso to be present, e.g., cyanide, several of which could adsorb to one Cd or evento a CdS moiety. By reducing the concentration of thiourea in the bath, C–N im-purities in the CdS film could be reduced almost to zero [77]. Whatever the natureof the C–N impurity, much of it could be removed simply by dissolution in waterat 60°C [78]

It should be pointed out that this deposition was carried out for films ca. 50nm thick; the study was carried out with CdS window layers for solar cells inmind, which are usually thin. It is possible that much longer depositions result indifferent impurities. Thus the sparingly soluble cadmium carbonate andcyanamide will be converted to CdS if enough sulphide ion is formed with time(or, for the complex-decomposition mechanism, if enough adsorbed thiourea de-composes on the surface of the solid phases). Of course, longer time also meansmore thiourea decomposition products.

Another study found much smaller concentrations of oxygen in the films(�4 at.%) [79], and most of this oxygen was attributed to bound water. Although


the deposition rates were comparable for this study and that in Ref. 76, the depo-sition solutions and conditions were quite different.

A decrease in the O content, measured by XPS, on Ar ion sputtering to-gether with a Cd:S ratio close to unity led Danaher et al. to propose that the O (notquantified) was present as surface sulphate [38]. In this same investigation, SIMSanalysis (much more sensitive than XPS, which is limited to ca. 0.1% concentra-tion) found a variety of impurities, including decomposition products of thiourea,CdO, and Cd(OH)2, but these were not seen in XPS, showing that they were pre-sent in very low concentration.

It is worth noting that analysis of the deposition solution after depositionwas complete, and after filtration of the solid precipitate showed the presence ofurea and guanidine, but not cyanamide, and that the amounts of these compoundswere less than those stoichiometrically expected, suggesting further decomposi-tion of urea to ammonia and carbonate [75] (see Sec. 3.2.1.1).

Films deposited from chloride and iodide (otherwise standard) baths werecompared [80]. I (ca. 3 at.%) was found in the iodide-deposited films but �1% Clin the chloride-deposited ones. About 5 at.% O was also found in both films. Theexcess Cd was believed to occur as Cd–O, Cd–OH, and, additionally in the iodidefilms, Cd–I or Cd–(I–O) species.

In most cases, the Cd:S ratio in these films was slightly greater than unity(usually between 1.02 and 1.1). A ratio of less than unity (0.92) was found formultiple layers (i.e., two or more layers deposited one on the other [47]; for asingle layer using the same deposition solution, the ratio was unity). The oxy-gen concentration varied from 8% (from an iodide bath) to 10–12% (from achloride or sulphate bath). As in the previous study, the oxygen was believed tobe present mainly as Cd(OH)2 or CdO. Even larger concentrations of oxygenwere found at low concentrations of ammonia (up to 18%) or at lower deposi-tion temperature. Another (XPS) study, however, found the ratio to be typically1.3 [81],

Some general conclusions can be drawn concerning oxygen in the films, inspite of the large spread reported in different studies, both in amounts and in in-terpretation of its source. The first thing to note is that, since the crystal size of theCdS is often in the region of 10 nm, around 10% of all the atoms will be locatedat a crystal surface. Thus adsorption of either oxygen or water will already showa relatively large amount of oxygen. However, unless this oxygen substitutes forsulphur, such adsorption will not change the Cd:S ratio. Other sources of oxygen,such as hydroxide, carbonate, and oxide (the last is less likely) will increase thisratio. As discussed earlier, carbonate can form from either dissolution of atmo-spheric CO2 or from the decomposition products of thiourea. Cd(OH)2 is morelikely to be formed when the pH is high (unbuffered solutions) or the ammoniaconcentration is low (less complexation, which probably outweights the slightlylower pH). The effect of temperature on Cd(OH)2 formation is complicated. A


higher temperature means a higher hydroxide concentration at a constant pH andless effective complexation on one hand, but also faster decomposition of thiourearesulting in more efficient conversion of the hydroxide to sulphide on the other.From the results of Nakada et al. [81] discussed earlier, where more oxygen isfound at lower temperatures, it appears that the latter effect is dominant.

Rapid thermal annealing (RTA) in vacuum of CD CdS films has beenshown to remove most of the oxygen that occurs in these films [82]. Typical an-nealing conditions were: heating rate—100°C/min; maximum temperature—600°C for 1 min; cooling rate—50°C/min. X-ray photoelectron spectroscopy(XPS) showed that, apart from the immediate surface, oxygen was effectively re-moved from CdS films (deposited from a thiourea bath). Dark resistivity was dras-tically reduced after this treatment, from ca. 107 �-cm for the as-deposited film toca. 1–10 �-cm after the RTA treatment. As can be expected, the photoconductionsensitivity also decreased drastically, from a light:dark resistivity ratio of ca. 104

to ca. 1.6 after annealing. This decrease in resistivity was attributed to removal ofelectron traps that originated from the adsorbed oxygen. The authors also sug-gested that, unlike conventional annealing, which results in a loss of stoichiome-try, e.g., by formation of Cd vacancies if annealed in vacuum, the RTA processdoes not change stoichiometry.

Pronounced thickness effects on resistivity were also noted in these RTACdS films. For example, a 95-nm-thick film showed a resistivity of 15 �-cm,which decreased to 0.2 �-cm for a thickness of 150 nm (with no pronouncedchange for even thicker films). Also, storage (in a dessicator, presumably in air)increased the resistance of the thin films, about an order of magnitude for the 95nm film after 50 days, with continuing increase but, for a 250 nm-thick film, onlya small, initial increase (ca. 50%). RTA of the stored films decreased the resistiv-ity to their original value before storage. The effect of storage was attributed tooxygen adsorption.

4.2 CdSe

4.2.1 A Mechanistic Introduction

Before going into details of the various aspects of specific CdSe depositions, andalthough it is not intended to deal with mechanistic aspects here (they have beenconsidered already in Chap. 3), it bears mentioning that, although in contrast toCdS, the complex-decomposition mechanism has not been discussed with respectto CdSe deposition, it is still possible that this mechanism does occur in some, oreven many, cases. If there is no evidence specifically in favor of this mechanismin general, there is also none against it. This point is stressed here since, in the lit-erature on CdSe (and selenides in general), it is automatically assumed that the re-action proceeds via free selenide ions.


4.2.2 Selenosulphate as Se Source

CdSe was deposited using a selenosulphate source as far back as 1970 [83]. Mir-ror films of CdSe were reported to form from an ammoniacal solution of Cd2�

only under conditions where Cd(OH)2 was present in the solution and at a pH �11.75. Under the conditions of those experiments, it was reported that this was theminimum pH required to convert Cd(OH)2 to CdSe. At higher values of pH, therate of conversion to CdSe increased, but so did homogeneous precipitation, withthe result that the films were thinner.

A detailed study of CdSe deposition was carried out using an ammonia-complexed solution with selenosulphate [84]. Most of this study was concernedwith kinetic measurements, already discussed in Chapter 3. Two different types ofsolution were considered: a clear solution where there was no visible Cd(OH)2 andone with added KOH to give a visible Cd(OH)2 suspension. The former requiredheating to at least 45°C for deposition to occur (although it is likely that deposi-tion would occur even at room temperature after enough time). The CdSe was ofthe zincblende structure. With a visible Cd(OH)2 suspension present, depositionoccured at room temperature, but the terminal thickness was only ca. 80 nm. Thehigher the pH, the lower the terminal thickness, since more of the Cd precipitatedin the solution. The CdSe from this bath was a mixture of wurtzite and zincblendestructures. The deposition rate and terminal thickness of the films were somewhatdependent on the nature of the substrate, both somewhat larger for Ge and Si thanfor glass.

A modification of this method used lower concentrations of ammonia (0.2M for a Cd concentration of ca. 50 mM) in a sealed vessel, thus preventing irre-producibilities due to escape of ammonia vapor [85]. Treatment of the Ti andstainless steel substrates by soaking in a suspension of Cd(OH)2 improved the ho-mogeneity of the films. At an ammnonia concentration of 0.3 M, no depositionoccurred (at least within the time frame of these depositions—about one hour).

A triethanolamine/ammonia bath has been used for CdSe [19]. While thissystem resulted in thick films for CdS (up to a few microns), CdSe films depositedunder the same conditions, only using selenosulphate instead of thiourea, werethinner [although films of 500 nm were obtained at 30°C that did not show signsof satuation (of thickness) after 25 hr—the longest time measured]. Ethylenedi-amine has also been used as a complex for Cd, with both precipitates and filmsformed [86]. In this case, the emphasis was on the precipitates, and no character-ization was carried out on the films.

Nitrilotriacetic acid (NTA) [N(CH2COOH)3] is a complexant for manymetal ions (see Sec. 2.9.1.2 for information on this compound). The sodium orpotassium salts of NTA have been used to complex Cd for CD of CdSe from se-lenosulphate solutions [10,87]. The rate of film growth depends on many factors,as discussed in Chapter 3; experimental details for CdSe deposition from this so-lution are given in Chapter 2. However, growth times are generally longer than for


ammonia-based baths—typically from a few hours to a few days for film thick-nesses in the range of 100—300 nm. The most notable property of these films istheir change in color with deposition conditions—from yellow if deposited at tem-peratures below 0°C and under conditions where the hydroxide mechanism is op-erative, to very deep red (thick films appear black by reflected light) for high-tem-perature, ion-by-ion depositions. This variation in color is a consequence of sizequantization, discussed in detail for these films in Chapter 10, and the CdSe (al-ways zincblende) crystal size varies from 3 to 20 nm. This color variation can betranslated into a variation of CdSe bandgap from ca. 2.3 eV for the yellow filmsto the bulk value (for zincblende CdSe) of ca. 1.8 ev. Annealing the films causescrystal growth and therefore loss of the size quantization effects. The major crys-tal growth, which corresponds to the phase change from zincblende to wurtzite,occurs between 300 and 400°C [88,89].

Since the optical spectra of these films are so sensitive to the crystal size (achange of 10% in crystal size can result in an easily measured spectral shift), mea-surements of the spectra provide a sensitive technique to investigate the effect ofdifferent deposition parameters on crystal size. Thus, while the crystal size is notstrongly dependent on the various concentrations of reactants (apart from theNTA:Cd ratio in the region where the mechanism changes), small increases (of theorder of 10–20%) in crystal size are observed if the Cd and/or selenosulphate con-centration is decreased considerably in the hydroxide cluster mechanism regime[90]. This can be rationalized, in a general way, by the greater likelihood of small,thermodynamically unstable CdSe nuclei growing to a stable configuration if thereactant concentrations are greater (see Section 1.2 for a discussion of nucleationand growth), since the growth rate will be faster. The greater the concentration ofnuclei, the smaller the final crystal size for a fixed reactant concentration. Thisreasoning can also explain the observation that the average crystal size increasessomewhat as deposition proceeds and reactant concentration decreases [89,91], al-though this growth might also occur by deposition of new CdSe on previously de-posited crystals; probably both mechanisms are operative to a greater or lesser ex-tent. From broadening of photoluminescence peaks with increasing depositiontime, it was inferred that the crystal size distribution also increased with deposi-tion time [91].

Illumination of the solution during deposition by supra-bandgap light af-fects crystal growth, probably via photoelectrodeposition of CdSe on the growingcrystals [90,92]. The solution used for CD can also be used for electrodepositionof CdSe. Light absorbed by the individual CdSe crystals forms electron/hole pairs,and the electrons can reduce the solution at the CdSe surface in the same way asthose supplied by an external power supply. Figure 4.3 shows the transmissionspectra of CD films deposited in the dark and under illumination for two differentdeposition temperatures. The red shift in the spectra of the illuminated films indi-cates a larger crystal size of these size-quantized samples (by ca. 1 nm) comparedto the nonilluminated ones. Additionally, the shape of the spectrum changes for


the low-temperature illuminated film; the onset becomes less steep. This spectralshape is typical for films electrodeposited from this solution [87]. Since the rateof electrodeposition is essentially temperature independent while that of CD isstrongly temperature dependent, the effect of the illumination (through the rela-tive amount of photoelectrodeposited CdSe) will be greater for low-temperaturefilms, seen particularly clearly by the pronounced change in shape of the low-tem-perature film. The growth does not occur for very weak illumination, suggestingthat the photoelectrochemical deposition is not very efficient, and other processes(electron/hole recombination or parasitic electrochemical reactions) dominate.Also, the crystal size saturates for light intensities above a certain level. This wasinterpreted to mean that one electron/hole pair was sufficient to influence thegrowth process [92]. This would not be expected for a photoelectrochemicalgrowth process, as described earlier. A more logical explanation for the saturation,particularly in view of the probable low quantum efficiency of the photoelectro-chemical CdSe deposition discussed earlier, is that one of the charges is removedmore rapidly into the solution than the other. If, e.g., this is an electron, then a holewill be left (probably trapped) on the crystal. Absorption of another photon willform another electron/hole pair, which will then recombine rapidly by Auger re-combination (three-body interaction) rather than form more CdSe.

While the pH of the deposition solution (based on the cluster mechanism)has been found to increase by as much as 0.8 during the deposition (see Sec.3.3.2), this increase was found to be considerably greater, up to 2.2 pH units, un-der illumination [92]. This could provide some clues about the mechanism of theCdSe formation under illumination. A possible pathway that could account for

FIG. 4.3 Transmission spectra of CdSe films deposited from a selenosulphate/NTA bathin the dark and under illumination (tungsten halogen lamp) at two different temperatures(6°C and 55°C).


this increase in pH (i.e., overall generation of OH�) is given as

Cd(OH)2 � SeSO32� � 2e� → CdSe � SO3

2� � 2OH� (4.5)

This reduction has, of course, to be balanced by the local oxidation reaction. Themost likely reaction, oxidation of sulphite, leads to an increase in acidity thatwould cancel out the rise in pH of Eq. (4.5). Probably other oxidation reactionsthat do not generate acidity occur that result in a net increase in pH (an examplewould be photocorrosion of the CdSe by photogenerated holes).

Another possibility that could explain the effect of illumination is a changein the electric double layer surrounding the CdSe particles, either adsorbed on thesubstrate or in the solution, which could lower a potential barrier to adsorption andcoalescence, as suggested previously for film formation from Se colloids under il-lumination [93]. Partial coalescence would reduce the blue spectral shift due tosize quantization. However, the spectral shape is not expected to undergo a fun-damental change in this case. The photoelectrochemical explanation therefore ap-pears more reasonable.

Addition of silicotungstic acid (STA) to a selenosulphate/ammonia/tri-ethanolamine bath resulted in a reduced rate of deposition but a larger final thick-ness (greater than 1 �m could be obtained). This was attributed to adsorption ofthe STA on the individual CdSe crystals (typically 4–5 nm in size), which impedesaggregation of the invidual crystals [94]. Reduced aggregation will slow downboth film growth (which relies on aggregation if the mechanism is one of colloidalgrowth as appears to be the case here) and precipitation of CdSe in solution, whichwill result in loss of CdSe for film formation. From the XRD pattern of such CdSefilms reported in an earlier study [95], a crystal size of ca. 4 nm could be esti-mated. The STA may also play a role in limiting the crystal size by capping, al-though even without the STA, small crystal sizes ca. 5 nm are usually obtainedfrom similar deposition baths.

4.2.3 Selenourea Source

Films have been deposited using selenourea and an ammonia-complexed solution at65°C [96]. Zincblende CdSe was obtained with an optical spectrum correspondingto a bandgap of 1.84 eV (the bulk room-temperature bandgap of zincblende CdSe isca. 1.8 eV). Analysis of electrical conductivity measurements indicated charge trans-fer occurred via a variable hopping mechanism through fairly deep states (a level0.29 eV below the conduction band was found from these measurements).

4.2.4 N,N-Dimethylselenourea Source

N,N-Dimethylselenourea was used with ammonia baths and additional citrate ortartrate complexation at pH of 11.3 (citrate bath) or 10.4 (tartrate) to deposit CdSeon glass at room temperature [97]. No XRD of the films was detected. From the


optical spectra, the bandgaps were high (ca. 2.1 eV, even larger for thinner films)suggesting crystal sizes of ca. 5 nm or less (see Chap. 10). The XRD spectra ofsuch a small crystal size can be missed in normal powder XRD if special care isnot taken, mainly slow scanning. Electrical and photoconductive properties ofthese films are described in Section 4.2.7.2.

Another study of these films concentrated on the particle size of the filmsand is discussed in Chapter 10 [98].

4.2.5 Selenosemicarbazide Source

One example is given in the literature where selenosemicarbazide was used to de-posit CdSe from solutions containing different complexants [99]. This reagentwas apparently more stable than selenourea or N,N-dimethylselenourea, and it issurprising that it does not appear to have been subsequently used. The CdSe filmswere specular and had a resistivity of 108–109 �-cm, which dropped about an or-der of magnitude on illumination.

4.2.6 Epitaxial Deposition

Using a nitrilotriacetate solution and a complex:Cd ratio high enough to preventCd(OH)2 formation (ion-by-ion mechanism), epitaxial growth of zincblende CdSewas obtained from a selenosulphate solution on both (111) and (1̄1̄1̄) polar facesof single-crystal InP [100]. (For lower complex:Cd ratios, in the regime of the hy-droxide cluster mechanism, the deposits were always polycrystalline, as expectedfor this mechanism.) The degree of epitaxy improved with increasing temperatureand was high at 90°C. Additionally, there was a high density of twins in the de-posits obtained at low temperatures, but less in those obtained at 90°C. The addi-tion of silicotungstic acid to the deposition solution destroyed the epitaxialgrowth, presumably due to blocking of the InP surface (and also the growingCdSe) by the strongly adsorbed silicotungstic acid.

This study also reported that films deposited on carbon membranes at tem-peratures �80°C were of hexagonal (wurtzite) structure, with a high density ofplanar defects, in contrast to the zincblende obtained from both hydroxide and ion-by-ion mechanisms at lower temperatures and to the epitaxial films on InP at alltemperatures.

4.2.7 Some Specific Optoelectronic Properties

4.2.7.1 Photoluminescence (see also Sec. 10.2.3)

Photoluminescence of CdSe films deposited from the selenosulphate/nitrilotriac-etate bath varied both in intensity and in spectral shape. As an example of the for-mer, a sample on glass, broken from a glass slide when the film thickness was ca.30 nm, gave a much stronger signal than the original sample left in the deposition


solution until the thickness reached ca. 100 nm. Such an effect, however, was notreproducible (unpublished results). Regarding spectral shape, both bandgap emis-sion with weak, if any, subbandgap signal [87,91] and deep subbandgap signalwith relatively weak bandgap emission [89,101] have been observed with no ap-parent difference between the samples. (Note that the use of the term bandgapemission refers to emission close to the bandgap energy. The emission may befrom very shallow traps.) Although the measurement temperatures for these dif-ferent experiments were not all the same, temperature does not usually have a verypronounced effect on the ratio between the two peaks, except below 50 K, andcannot explain the observed differences. In humid atmosphere, where water vaporis adsorbed on the CdSe, the predominant emission is close to bandgap [101] (seeSec. 4.2.7.3). Most studies do not state whether measurements were carried out inthe open atmosphere (low-temperature measurements, of course, are not). An in-vestigation on films deposited from an N,N-dimethylselenourea/citrate/ammoniabath showed both bandgap and lower-energy (ca. 1.75 eV) peaks [98]. The latterwas attributed to larger crystal size (see Chap. 10), but it is likely that the low-en-ergy peak is a subbandgap response arising from surface states.

The temperature dependence of the emission spectra provides useful infor-mation on the source of the emission. Figure 4.4 shows a series of emission spec-

FIG. 4.4 Photoluminescence spectra taken at different temperatures of a CdSe filmdeposited from a selenosulphate/NTA bath. Crystal size of CdSe ca. 4 nm. (From ref. 101with permission from Elsevier Science).


tra of a (4-nm-crystal-size) CdSe film taken at different temperatures. Each spec-trum is made up of a bandgap peak (at high photon energy) and a subbandgap peakat lower energy. The bandgap peak becomes weaker as the measurement temper-ature increased. This is normal for a band-to-band transition. The subbandgappeak, however, initially increases up to ca. 50 K, and then decreases with furtherincrease in temperature. This behavior (and also incident light intensity depen-dence of the emission) is typical of donor–acceptor recombination. In the presentcase, the nanocrystals are essentially intrinsic and are not expected to contain bulkdopants. However, the surface states (see following section), after trappingcharge, can behave in much the same way, and the recombination can be ex-plained by recombination of surface-trapped electrons and holes with the emissionred-shifted ca. 0.5 eV from the bandgap [101]. In the study by Trojanek et al. [89],this emission is shifted nearly 0.7 eV from the bandgap and the parallel increasein emission energy as a function of crystal size with the theoretical (effective massapproximation) conduction band shift interpreted to mean that the emission oc-curred from shallow-trapped electrons to deep-trapped holes.

Time-resolved measurements of photogenerated (very intense illumination,up to 0.56 GW/cm2) electron/hole recombination on CD (selenosulphate/NTAbath) CdSe of different crystal sizes has shown that the trapping of electrons, prob-ably in surface states, occurs in ca. 0.5 ps, and a combination of (intensity-depen-dent) Auger recombination and shallow-trapped recombination occurs in a timeframe of ca. 50 ps. A much slower (not measured) decay due to deeply trappedcharges also occurred [102]. A different time-resolved photoluminescence studyon similar films attributed emission to recombination from localized states [103].In particular, the large difference in luminescence efficiency and lifetime betweensamples annealed in air and in vacuum evidenced the surface nature of these states.

A photoluminescence study of CdSe deposited from a selenourea/ammoniasolution onto glass at 80°C and relatively low pH (7–8) was made [31]. An emis-sion peak centered at 1.445 eV (860 nm) was observed with a tail to the low en-ergy side. Such an emission must be due to deep traps, since the shift from thebandgap emission is ca. 0.4 eV, a value close to, or somewhat less than, that forfilms deposited from selenosulphate solution (see earlier). Annealing in air shiftedthe emission to higher energies.

4.2.7.2 Photoconductivity

There appears to be a fundamental difference between films deposited from a se-lenosulphate source and those deposited from a dimethylselenourea source (mostof the detailed photoconductivity studies are on the latter). As-deposited films us-ing a dimethylselenourea source had a resistivity of ca. 2 � 1012 �/sq (�108

�-cm for a film thickness of 0.5 �m). The resistivity dropped only a little underillumination, with a maximum decrease of less than an order of magnitude. For


films deposited at a higher temperature (50°C), the resistivity was up to an orderof magnitude lower (ca. 107 �-cm), which dropped ca. six times under illumina-tion. Air-annealing increased the photosensitivity greatly (up to ca. 107 at 450°C),in contrast to the case for CdS, where good photosensitivity was obtained for as-deposited films (but always deposited at high temperatures) and was reduced onair-annealing. The dark resistivity of the CdSe films increased by nearly an orderof magnitude after annealing, again in contrast to the decrease usually obtained forCdS films. This behavior is also in contrast to CdSe films deposited from seleno-sulphate/triethanolamine/ammonia solutions, where the relatively low resistivityof the as-deposited film (5�103 �-cm) dropped to a few ohm-centimeters on annealing (430°C in air) [95], or to another study of selenosulphate/ammoniafilms with dark resistivity of 107–108 �-cm, which dropped to 1–10 �-cm afterannealing at 280°C in vacuum [104]. On the other hand, in another study on se-lenosulphate/ammonia films, which reported a dark resistivity of 105–106 �-cmas deposited, the dark resistivity initially increased on heating in air (ca. four timesat 180°C) then decreased only an order of magnitude on heating at 340°C [105].The photosensitivity of these latter films was low but appreciable in the as-de-posited state (varying from 5 to 50) and increased up to ca. 103 after annealing at180°C before decreasing again at higher annealing temperatures. While it is clearthat CdSe films deposited from dimethylselenourea possess higher photosensitiv-ity after annealing compared to those deposited from a selenosulphate bath, thereasons for the difference are not understood. One possibility that might explainthis difference was suggested, based on a comparison of CdSe films depositedfrom a dimethylselenourea bath with CdS films. It was hypothesized that the dif-ference in the effect of annealing on dark resistivity was due to the formation ofconducting CdO due to the oxidation of CdS, which did not occur in CdSe [106].The ease of oxidation of the Cd chalcogenides is normally CdTe � CdSe � CdS;however, this usually refers to oxidation of the chalcogen, e.g., CdTe to CdTeOx,rather than to CdO. More details on the effect of annealing on the photoconduc-tion properties are given in Refs. 106 and 107. Another characteristic of thesefilms, different from most CD CdS films is the fast photocurrent decay (no morethan a few seconds at most) of the films, both as deposited and annealed. Clearly,the trapping centers in these films are also efficient recombination centers, whichmay, at least in part, explain the low photosensitivity of the as-deposited films.CdSe films prepared by the selenosulphate/citrate process were less photosensi-tive than those deposited by the dimethylselenourea method, although details werenot given [106].

Some studies of these films were made after immersion in solutions of Hgor Cu ions, when ion exchange reactions occurred, converting the surface of thecrystallites to (partial) Hg–Se and Cu–Se compounds. As expected, such treat-ments could affect the photoconductivity of the films greatly.


4.2.7.3 Surface States on CdSe Films

An important aspect of semiconductor films in general with regard to electronicproperties is the effect of intrabandgap states, and particularly surface states, onthese properties. Surface states are electronic states in the “forbidden” gap that ex-ist because the perfect periodicity of the semiconductor crystal, on which bandtheory is based, is broken at the surface. Change of chemistry due to bonding ofvarious adsorbates at the surface is often an important factor in this respect. ForCD semiconductor films, which are usually nanocrystalline, the surface-to-volume ratio may be very high (several tens of percent of all the atoms may be sit-uated at the surface for 5 nm crystals), and the effects of such surface states are ex-pected to be particularly high. Some aspects of surface states probed by photolu-minescence studies are discussed in the previous section.

Surface treatments of CD CdSe films deposited from selenosulphate/NTAsolutions have a pronounced effect on various optical, electrical, and optoelec-tronic properties of the films, due to interaction with or modification of such sur-face states. Mild etching (dilute HCl) of the films reverses the direction of currentflow both in CdSe/polysulphide photoelectrochemical cells [108] and in Kelvinprobe surface photovoltage (SPV) measurements in air [109]. These studies arediscussed in more detail in Chapter 9, in Section 9.2 on photoelectrochemicalcells. At this point, it is sufficient to state that the effect is believed to be due topreferential trapping of either electrons or holes at surface states that are modifiedby the etching process.

The adsorption of water vapor on these CdSe films acts to passivate, at leastto a large extent, some of these surface states. In particular, strong subbandgap sig-nals have been observed in SPV measurements of these CdSe films (as deposited)only when measured in a dry ambient; in normal atmospheric humidity, no suchsignal occurs and only suprabandgap light gives rise to an SPV signal [109]. Par-allel results have been observed in photoluminescence measurements, which areparticularly sensitive to surface states: The predominant subbandgap emission thatoccurs in a dry atmosphere changes to a predominantly (near) bandgap emissionin a humid atmosphere [101]. The asymmetrical nature of these states, seen in op-tically detected magnetic resonance (ODMR) spectroscopy, is further evidencefor their surface nature; bulk states are expected to be symmetric [110]. It is im-portant to note that these effects are seen only in small-crystal-size nanocrystallinefilms (the foregoing experiments were carried out on 4- to 5-nm-crystal-sizefilms). No such effects were observed if the crystal size was ca. 20 nm; the sur-face-to-volume ratio is already much smaller for this size. Current–voltage spec-troscopy of individual CdSe quantum dots deposited mostly by electrodeposition,but also by CD, using a conducting AFM (atomic force microscope) tip alsoshowed directly the presence of surface states in a dry atmosphere but not in a hu-mid ambient [111].


4.3 CdTe

There appear to be only two independent reports in the literature on CD CdTe(which is the only telluride reported in the CD literature). In the first one, the CdTewas deposited with the main purpose of carrying out electron spin resonance andmorphological studies of the effect of annealing, while another paper, based ap-parently on the first, described the deposition, with the main purpose of further usein photovoltaic cells. Therefore only limited detail on the actual deposition orproperties of the as-deposited films were given. This is surprising in view of thefact that the first report is the first case of a CD telluride.

In the original study by Padam and Gupta [112], the deposition solution con-tained triethanolamine/ammonia-complexed CdAc2 and the Te source was TeO2

with hydrazine hydrate as a reducing agent. The nature of the TeO2 solution wasnot clear, since TeO2 is only slightly soluble in water; it may have been dissolvedin a hydroxide solution, in which it is much more soluble. The deposition was car-ried out at 90°C. Electron diffraction of the as-deposited films showed bothzincblende and wurtzite phases of CdTe (the zincblende phase is the more stableand commonly encountered one).

Buckley used the same technique to deposit films for photovoltaic cells[113], only with CdCl2 as Cd source and apparently a lower Cd concentration.Electron diffraction of these films showed a predominantly zincblende structurewith some wurtzite phase. The films were p-type with a resistivity of 20 � 5�-cm. These values, and the subsequent photovoltaic cells, apparently refer to as-deposited films; no reference to annealing was made in this study.

The second, more recent investigation described Te dissolved in sodiumsulphite as a source of telluride (tellurosulphate) [114]. It was previously believedthat Te was insoluble in sulphite solution under normal conditions, although thereis one previous reference in the literature to this reagent prepared under pressureat high temperatures [115]. The CdTe deposition described in Ref. 114 indicatesthat the solubility is sufficient to allow deposition of tellurides. The depositionwas carried out in a solution of CdSO4 containing triethanolamine, ammonia, andNaOH. Both deposition rate and film thickness were maximal at 75°C depositiontemperature. As with the previous TeO2 deposition, both zincblende and wurtzite(dominant) phases of CdTe were obtained. Elemental analysis showed a small Cdexcess. This appears to be in contradiction to the XRD analyses, which showedconsiderable amounts of Te (also TeO2, particularly after heating at 100°C). Al-though the films were apparently highly scattering, making bandgap measurementmore difficult, the bandgap (direct), measured from the optical spectrum, was ca.1.4 eV, close to the literature value. The (room-temperature) resistivity was ca.106 �-cm and the conductivity n-type. The carrier density was ca. 1019 cm�3.

In experiments carried out by us to repeat these two methods, we have beenable to deposit CdTe, although stoichiometry control was difficult, with Te oc-


curring in the hydrazine method and mainly TeO2 impurity in the tellurosulphateone. Probably careful attention to removal of oxygen from the solution and maybeaddition of hydrazine to the tellurosulphate-based solution would improve the sto-ichiometry. It is clear that other telluride films, in particular with metals that formvery insoluble chalcogenides, such as Pb, Cu, and Hg, should be accessible usingthese methods.

Although not strictly CD, CdTe films were very recently prepared by treat-ing CD Cd(OH)2 films with a telluriding solution [116]. The Cd(OH)2 films weredeposited from an alkaline H2O2 bath containing citrate-complexed Cd. Thesefilms were treated with a solution of elemental tellurium dissolved in hydrox-ymethanesulphinic acid. It appears that this solution contains telluride ions, al-though it has not yet been well characterized. The Cd(OH)2 was converted (incompletely) into CdTe films. The bandgap was ca. 1.5 eV (approximate due tothe highly scattering nature of the films). The dark resistivity was ca. 5 � 108

�-cm, which decreased to ca. 7 � 107 �-cm upon illumination.

4.4 ZnS

4.4.1 Introduction

Chemical deposition of ZnS has been the subject of considerable activity, the mainreason for which is its hoped-for substitution for CdS in thin-film photovoltaiccells. Since the chemistries of Zn and Cd are similar in many ways, it might be ex-pected that deposition of their chalcogenides is also similar. However, there is adominant difference in their properties that results in the fact that ZnS is consider-ably more difficult to deposit by CD than CdS. This difference is manifested by thedifference in solubility products between the respective hydroxides and chalco-genides. Considering, for example, the sulphides, the relevant values of Ksp are:

Cd(OH)2 2 � 10�14 CdS 10�28

Zn(OH)2 8 � 10�17 ZnS 3 � 10�25

The deposition for mechanisms proceeding through hydroxide clusters is depen-dent on a large difference between the solubility products of the hydroxide andsulphide, since the sulphide exchanges the hydroxide. The situation for ZnS istherefore much less favorable than that for CdS: About a million times higher con-centration of sulphide is required to form ZnS than CdS [see Eq. (3.46)], as shownin graphical form in Figure 4.5. More sulphide is also required at higher tempera-tures because of the strongly increasing ion product of water with increasing tem-perature, resulting in higher hydroxide concentrations for any particular value ofpH. (It is again stressed that in dealing with values of solubility products, there arelarge variations, sometimes of orders of magnitude, between one source and an-other. Therefore calculations based on these values are correspondingly impre-


cise. However, to take the present case as an example, whether the difference insolubility product ratios between the Zn and Cd hydroxides and sulphides is 106

or an order or so larger or smaller in magnitude does not qualitatively alter theconclusion.)

The obvious solution to this problem is to deposit ZnS at a lower pH whenthe OH� concentration will be lower. However, with thiourea, for example, low-ering the pH results in slower hydrolysis of the thiourea and therefore a lower sul-phide concentration (presumably also a reduced decomposition rate if thethiourea–hydroxide complex-decomposition mechanism is effective). This hasbeen circumvented, as described later, by working with alternative sulphidingagents at low values of pH.

Deposition by a pure ion-by-ion mechanism should also solve this problem,since no hydroxide is involved. However, in this case we encounter the problemof the high Ksp of ZnS compared to CdS, which again means that more sulphide isneeded. For thiourea, this means a higher pH, which again means that strong com-plexation is needed to prevent Zn(OH)2 formation, by reducing the free [Cd2�].However, this will also reduce the rate of ZnS deposition. While there are manyexamples in the literature of cluster deposition of ZnS, there does not even seemto be one unambiguous case of ion-by-ion deposition of this semiconductor.

As already implied, most depositions of ZnS have been carried out underconditions where Zn(OH)2 can be formed. From the forgoing general discussion,this means that, even if ZnS is identified (by XRD, for example), there is in moststudies no evidence for the absence of some hydroxide species. More recent stud-

FIG. 4.5 Steady state concentration of sulphide ion needed to convert the hydroxides ofCd and Zn into the corresponding sulphides at 25°C and 60°C.


ies have addressed this problem, and there are some papers with “Zn(OH,S)” inthe title. This probably usually refers to a mixture of ZnS and unreacted Zn(OH)2

and not to a true ternary (more accurately, quaternary) compound. For that reason,such films are discussed in this chapter rather than in Chapter 8 (Ternary Semi-conductors).

A recent condensed review on ZnS deposition, with emphasis on the differ-ences in CD of CdS and ZnS, is given in Ref. 117.

The majority of studies involved ammonia-complexed baths with thiourea asa source of S. Hydrazine was also used in most cases, although there were severalstudies where ZnS was obtained without hydrazine. Amines (triethanolamine orethanolamine) were also used, again both with and without hydrazine (which is it-self a type of amine). One effect of the hydrazine is to speed up the deposition. Hy-drazine, a strong reducing agent is expected to increase the rate of sulphide forma-tion. However, it seems that this is only part of the picture. A study of the effect ofvarious amines on the rate of ZnS deposition showed that although hydrazine gavethe fastest rate, other amines (ethanolamine and triethanolamine) also increased thedeposition rate [118] (see Fig. 3.7). The amines all act as complexants; thereforethey would reasonably be expected to reduce the deposition rate by reducing thefree-Cd2� concentration. Amines have reducing properties (redox potentials of hy-drazine, ethanolamine, and triethanolamine are �1.16, �0.56, and �0.46, respec-tively). It appears that the amines accelerate thiourea decomposition. However, themechanism of this effect is not yet clearly understood [119].

Kinetic studies have found a value for the activation energy of ZnS film for-mation from thiourea/ammonia-based baths of 5 kCal/mole (21 kJ/mole), too lowto be chemical reaction controlled (unlike CdS, which is thiourea-decompositioncontrolled) but not diffusion controlled, since stirring the deposition solution hasno effect [120,121]. This was explained by a rate-determining step of Zn-liganddissociation from a hydrazine complex. However, addition of an additional com-plexant into the solution will only act to reduce the free-Zn2� concentration, notonly due to the extra complexing power of the hydrazine, but also because of thestatistical factor that complexes comprising more than one ligand will be furtherstabilized by the fact that different combinations of the ligands can be found in thecomplexes [119]. There is clearly a fundamental difference in deposition mecha-nism of these ZnS depositions compared with that normally encountered withCdS. More details on the difference in activation energies of deposition for CdSand ZnS (ZnSe) are discussed in Section 3.5.

4.4.2 Specific Studies of ZnS Deposition

Some specific properties of ZnS films grown from thiourea/ammonia/hydrazinebaths, sometimes with added compounds, are given next.

Doña and Herrero noted that hydrazine was not essential for growth but thatit speeded up the growth rate (by a factor of ca. 3) and improved the homogeneity


and the specular reflectance (a consequence of better microhomogeneity) [120].No XRD pattern was observed for these films, indicating either highly disorderedand/or very small crystals. However, TEM/ED did show that the sphalerite phasewas formed with a crystal size of �10 nm. The stoichiometry measured by EDSwas Zn:S � 1:1. This implies that either pure ZnS or a hydroxy sulphide(Zn,OH,S) was formed.

Optical transmission of the films was a spectral average of ca. 85% beyondthe bandgap. The bandgap measured from the spectrum was 3.76 eV (literature forzincblende ZnS � 3.6 eV), suggesting size quantization, although the measuredcrystal (ca. 10 nm) is at least twice the size required for such an effect.

Resistivity was nearly 109 �-cm and independent of bath composition.Temperature dependence of resisitivity gave an activation energy of 0.95 eV,which was ascribed, based on previous studies, to an acceptor level above the va-lence band due to Cu impurity.

Using various amines added to the ammonia bath (in most cases with addedhydrazine), sphalerite ZnS films were obtained with a crystal size of ca. 3 nm [118].Rutherford Backscattering Spectroscopy (RBS) analyses showed that there wasabout twice as much Zn in the films as S. (More basic solution and more hydrazinegave more stoichiometric films). Extended X-ray Absorption Fine Structure (EX-AFS) and Fourier Transform Infra-red (FTIR) spectroscopy showed that the filmsdid not have Zn-O groups but rather Zn-OH ones [122] and that there is probablya mixture of ZnS and unreacted Zn(OH)2, quite likely as a ZnS shell around aZn(OH)2 core. Optical spectra gave a bandgap of ca. 3.85 eV, considerably blue-shifted from the bulk value of 3.6 eV, as expected from such small crystals.

Nucleation studies of ZnO and ZnS on glass and SnO2-glass from ammo-nia/thiourea baths (sometimes also with hydrazine) were carried out [123]. The de-position conditions, mainly pH, were varied. On glass, both ZnS and ZnO could bedeposited, depending on conditions. On SnO2-glass, however, only ZnO wasformed (a few percent S could be obtained at high pH). This suggested that a sur-face-activated mechanism was important for nucleation of ZnS and less so for ZnO.

Most depositions were carried out at high temperatures. One exception wasa room-temperature deposition from an ammonia/thiourea bath (no hydrazine)[124]. No structural characterization was made, but the optical spectra were con-sistent with ZnS. The spectra show high transmission and low absorption in thesuprabandgap region, along with low specular reflectance. Some samples, whichhad been deposited for a relatively long time, exhibited strong absorption (andsomewhat increased reflectance) in the IR region, which could be ascribed to freecarriers. The bandgap, at between 3.7 and 3.8 eV, was slightly higher than the bulkvalue for sphalerite ZnS. Some room-temperature depositions were also carried outfrom a thiourea/ammonia/ammonium sulphate/hydrazine bath. Again, high trans-mission in the suprabandgap region and a bandgap of 3.75 eV were measured.

A bath without ammonia, using triethanolamine/hydrazine/thiourea andadded NaOH to a pH ca. 12.3, was described [125]. Contrary to most other depo-


sitions, the ZnS had the wurtzite structure, measured by XRD. Optical spec-troscopy gave a bandgap of 3.68 eV Transmission above the bandgap was ca.30%, implying a large amount of scattering (specular reflection would not be ashigh as 70%). The resistivity was ca. 106 �-cm, and the films were n-type.

Thioacetamide has also been used to deposit ZnS. In this case, depositionscould be carried out under acid or neutral conditions as well as alkaline ones.Three different studies were carried out using thioacetamide. In one [126], twodifferent baths were investigated—one with and the other without ammonia (thelatter was probably slightly acidic and certainly had a considerably lower pH thanthe ammonia bath). Zincblende ZnS was obtained from both baths. The crystalsize (TEM) was 6–8 nm (ammonia bath at 90°C); from the ammonia-free bath, thecrystal size was not given, but from the sharpness of the ED pattern, it is probablylarger. The bandgap (temperature independent) for films deposited from the am-monia-free bath was 3.6 eV (� literature value). For films deposited from theammonia bath, values of bandgap up to 4.05 eV (5°C deposition), which dependedon deposition temperature (3.8 eV at 90°C; 3.95 eV at 30°C), were measured.Urea was used in the ammonia-free bath to improve adherence, although it is notknown why this affected adherence.

ZnS was grown from a thioacetamide bath using triethanolamine to com-plex the Zn2� and buffered by NH3/NH4

� to a pH of 10 [127]. Glass coated withthese ZnS films was found to be a good substrate for other CD semiconductors thatmight otherwise exhibit poor adhesion, and this was the purpose of this study. Itwas noted from optical transmission that for films no thicker than ca. 0.15 �m,scattering was negligible, but it became increasingly marked for thicker films.XPS depth-sputtered analysis of these films indicated that the ZnS–glass interfacewas not sharp, and it was suggested that Zn diffuses into the glass to some extent,explaining the good adhesion of the films [128]. A more complete characteriza-tion of these films was subsequently carried out [129]. No XRD pattern was foundfor the as-deposited films, implying amorphous or very small crystal size. Evenafter a short anneal at 500°C, a crystal dimension as low as 13 nm (depending oncrystal orientation) was measured by XRD, implying a much smaller size in theas-deposited film. In line with the crystal size, optical absorption showedbandgaps between 3.85 and 3.95 eV, higher than the bulk value, presumably dueto size quantization.

These films were not photoconductive as deposited but became so on air-an-nealing with an optimum annealing temperature of 388°C, when the photosensi-tivity increased to 104. The film resistivity decreased from ca. 5 � 107 �-cm (asdeposited) to ca. 104 �-cm after air-annealing at 400–500°C.

The third study used a solution of ZnCl2 and thioacetamide at a pH of 2.45[130]. Films were deposited on ITO/glass at 70°C. In spite of the acid conditions,which precluded formation of Zn(OH)2, a precipitate formed in the solution in


parallel to film formation. The crystal size for both film and precipitate was ca. 3nm. The energy of activation was measured for both the precipitate and the filmformation (usually, it is only measured for film formation) and found to be thesame for both, with a value of 9.5 kCal/mole (40 kJ/mole); this value is interme-diate between that usually found for CdS deposition, where the rate-determiningstep is a chemical reaction, and for ZnS deposited from an alkaline thiourea bath[120] (21 kJ/mole; see Sec. 3.5). Whether this difference is due to the expecteddifferent mechanisms (ion-by-ion probable for the acidic bath and cluster mecha-nism for the alkaline one) or to differences in thiourea and thioacetamide decom-position is unknown at present.

Thiosulphate has been used in an acid bath (pH 2–4) at 85°C [131]. Thefilms were poor (not uniform, powdery, and nonadherent). The only other charac-terizations given were the bandgap (3.4 eV) and resistivity (106–107 �-cm).

Thiosulphate has also been used to deposit ZnS using a photochemical re-action, in the same manner as used previously for CdS (see Sec. 4.1.6.7 for a de-scription of the proposed mechanism). In brief, UV light of wavelength shorterthan 300 nm (from a Hg lamp) forms hydrated electrons and elemental S fromthiosulphate solution, and this reacts with Zn2� to give ZnS [132]. The ZnS washighly nonstoichiometric with excess Zn. This is not surprising, since ZnS itselfis sensitive to UV light, with the formation of elemental Zn due to the strong re-ducing action of photogenerated electrons.

While referring to precipitation in solution and not CD of films, it is of in-terest to mention that N-allylthiourea was used as a sulphur source in a Zn2�/am-monia bath [133]. Pure ZnS (as detected by XRD) was obtained at a pH of ca. 11.0at 90°C. At lower pH values (and at 80°C), only ZnO was obtained; at higher val-ues, a mixture of ZnO and ZnS was formed. The ZnS was the wurtzite form in allcases. The ZnS crystal size was ca. 5 nm at pH � 11.0 and slightly smaller (ca. 4nm) at higher pH values, which gave a mixed phase of ZnS and ZnO. The ZnOcrystal size was much larger (ca. 200 nm). The ZnS fraction increased as the am-monia concentration (at constant pH) increased (lower [Zn2� ]) or as the pH de-creased (at constant ammonia concentration) from 13 to 11 (lower [OH�]). Suchexperiments should help in choosing optimum conditions for CD.

A variant of CD has been described where ZnS was precipitated as a gel byadding concentrated S2� to a concentrated solution of a Cd salt. The pH was thenreduced by HNO3 to a value between 5 and 7, when a semitransparent sol formed.Heating this sol between 100 and 200°C (in an autoclave) resulted in the forma-tion of zincblende ZnS films [134]. If the S2� was not in excess (twice the Cd con-centration), some ZnO, together with some wurtzite ZnS, also formed. Additionof a CuCl2 solution to the pH-adjusted sol and heating at 140°C resulted in Cu-doped ZnS (particle size 60 � 10 nm) that showed several photoluminescencebands [135].


4.5 ZnSe

4.5.1 Introduction

In contrast to CdSe, which in most cases was prepared using selenosulphate as aSe source, ZnSe was deposited in the majority of studies using selenourea. One re-port was given where ZnSe (among other selenides) was deposited from an am-moniacal solution at 20°C [99] using selenosemicarbazide as a Se source. Otherthan that the ZnSe films were specular in nature, no details on their propertieswere given. There are two reported studies where selenosulphate was used. In one[136], it was noted that hydrazine was essential to obtain film formation; for ZnS,while it was usually preferable to include hydrazine in the deposition solution, itwas apparently not essential and films could be made without it. This differenceprobably reflects the higher solubility of ZnSe compared to CdSe, the lower solu-bility of the hydroxide of Zn compared to that of Cd and either the faster forma-tion of selenide ion from selenourea compared to selenosulphate or a different de-position mechanism for the different Se sources. The other study reportedlyobtained ZnSe without using hydrazine (triethanolamine and ammonia were usedtogether as complexants) [137]. Unusually high-concentration selenosulphate so-lution (0.8 M) was used. No structural or analytical characterization was carriedout on the films, although the bandgap measured from the optical absorption wasin the correct region for ZnSe.

Hydrazine was also used in all the studies employing selenourea (but notwhere N,N-dimethylselenourea was used—see later). It is not clear to what degreehydrazine was essential in these studies. In an early theoretical study of ZnSe de-position from Zn–ammine/selenourea baths, the use of hydrazine is not mentioned[138]. On the other hand, in a patent by the same authors describing ZnSe deposi-tion, the use of hydrazine hydrate was the main issue in the claims [138a]. In mostof these studies, ammonia was used as a complexant. The conditions varied some-what from study to study, but it appears, either explicitly or by resorting to edu-cated guesses, that in all cases the deposition occurred by a hydroxide clustermechanism. The pH was usually ca. 11.5, and deposition temperatures variedfrom 50°C to 80°C. The exception was the deposition described in the patent byKitaev and Sokolova, where only ZnCl2, hydrazine hydrate, and selenourea (�acetic acid to acidify the very alkaline solution to a pH of ca. 9) were used [138a].It was claimed that deposition could be obtained over a temperature range of10–70°C. No characterization or properties of these particular films were given.

4.5.2 Depositions Using Selenourea

4.5.2.1 Structure

XRD of the films gave very broad, ill-defined peaks [139,140]. In one study, theprecipitated powder from the deposition solution was found to be wurtzite, while


grazing-angle XRD (more sensitive for thin films) gave a poorly defined, verybroad spectrum [141]. In another study [140], while no normal XRD pattern wasobtained, grazing-angle XRD showed a mixture of wurtzite and zincblende ZnSetogether with a little ZnO. TEM/ED showed zincblende ZnSe in one study [139]and predominantly wurtzite in another [140]. Crystal size was measured in twocases by TEM: 2–2.5 nm [139] and ca. 10 nm [140].

4.5.2.2 Composition

Elemental analyses found such films to be Zn rich (with respect to Se) [139,142].The films were oxygen rich, and it is probable that, as was often found with ZnS,the films are a mixture of ZnSe and Zn(OH)2, together with a little ZnO in somecases [140]. An XPS depth profile study found the surface of the films to be morestoichiometric (although still Zn rich) than near the substrate [143].

4.5.2.3 Optical Spectroscopy (Bandgap)

Optical spectroscopy was used to measure bandgap (bulk value ca. 2.6 eV). Suchvalues could be correlated with the crystal size, with values of 2.7 eV [140] and 2.9eV [139,142], the latter due to size quantization in very small crystals (ca. 2.5 nm).One study noted that the film was a transparent white (colorless) turning to orange,presumably due to Se formation on exposure of excess selenourea to air [141]. Thewhite film, when annealed in air at 300°C, showed a broad ZnSe diffraction peak(by 400°C, the film was converted to ZnO). This, together with the very broadpeaks (of both the 300°C annealed sample and of the powder precipitated in the so-lution) suggested that the white color was due to size quantization (white implies abandgap of 3 eV—bulk ZnSe is pale yellow). The presence of oxide/hydroxidein these films would result in a weakening of the yellow color of pure ZnSe butwould not be expected to change the spectral position.

4.5.2.4 Electrical Resistivity

Electrical resistivity was measured in only one case, with a value of 2 � 108

�-cm [140].

4.5.3 Deposition Using N,N-Dimethylselenourea

A modification of the general procedure used a substituted selenourea (N,N-dimethylselenourea) instead of selenourea and no hydrazine. Citrate was used, to-gether with ammonia, presumably as a co-complexant, although it may also havefunctioned as a mild reducing agent, and the deposition was carried out at a rela-tively low temperature of 50°C [144].

No XRD pattern of the films was found, but the powder precipitated in thesolution was zincblende phase with very broad peaks (�3 nm coherence length).In contrast to other measurements, where the Zn:Se ratio was greater than 1, the


Zn:Se ratio of these films was less than 1 (0.85). As with the other films, therewere considerable amounts of oxygen (this was measured by XPS and includesnormal surface-adsorbed oxygen). A value for the bandgap of 2.63 eV was mea-sured from the optical spectrum. This (very slightly) higher value than the bulkvalue was explained by the small crystal size. (Size quantization for a crystal sizeof 3 nm would be expected to give a larger blue shift than this). The films were n-type with a resistivity of ca. 3 � 107 �-cm, within an order of magnitude of thevalue measured from a selenourea bath [140].

4.5.4 Deposition Using Selenosulphate

There is one reported brief study of ZnSe deposition using selenosulphate [136].Considering the understandable preference for using selenosulphate rather thanselenourea for CdSe depositions in most cases (selenosulphate is more stable andsimpler to make and to handle), it is surprising that this is not also the case forZnSe. It is possible that this is simply a case of “inertia”; i.e., most researchers follow essentially the same recipe (although the selenosulphate technique de-scribed here predated the other studies).

Hydrazine was used (in its absence, no ZnSe was formed). A mixture of bothtriethanolamine and ammonia, together with NaOH, was used as complexant/pHadjuster; no explanation of the solution composition was given. The deposition wascarried out at 100°C (the highest deposition temperature of all these processes).

The films were light yellow (characteristic of ZnSe), which changed to alight reddish color. The reddish color might have been due to formation of Se un-der the conditions of the deposition, although Se would be expected to dissolve inthe selenosulphate. X-ray diffraction showed the films (presumably the reddishones) to be wurtzite ZnSe. A bandgap of 2.62 eV was measured from the opticalspectra. The films were n-type with a relatively low resistivity of ca. 104 �-cm.

4.5.5 Deposition Using Selenosemicarbazide

Selenosemicarbazide was used in one study to make, among other materials, ZnSefilms and precipitates from aqueous ammoniacal solutions at 20°C [99]. The filmswere specular, but no further information was provided.

4.5.6 Miscellaneous Methods

A novel CD technique used metallic Al to reduce Se. Elemental Se and metallicAl foil, together with ZnCl2, were dissolved in NaOH solution and heated to 80°Cin an autoclave with substrates of Teflon or alumina [145]. (Metal chalcogenidefilms have been chemically deposited onto Teflon in a number of reports.) It is no-table that the films deposited onto alumina were reported to be 1.8 �m thick—much thicker than ZnSe CD from other solutions. The films were sphalerite ZnSe


with a crystal size greater than 10 nm, estimated from XRD peak widths. TheZn:Se ratio was essentially unity, in contrast to other reports, where the films wereinvariably nonstoichiometric. (At deposition temperatures of 125°C and above,ZnO was also formed.) The mechanism of this process was suggested to be:

3Se � 6OH� → 2Se2� � SeO32� � 3H2O (4.6)

2Al � SeO32� � 3H2O � 2OH� → 2Al(OH)4

� � Se2� (4.7)

Se2� � Zn(OH)42� → ZnSe � 4OH� (4.8)

Since the presence of the Al was essential to obtain films, Reaction (47) or otherreactions, possibly involving reduction of Se species with nascent hydrogen(formed by the dissolution of Al in the NaOH solution), were probably the im-portant steps in the deposition.

Another method of ZnSe deposition is not true CD but is related and worthyof mention. H2Se, prepared by decomposition of CdSe with HCl, was passed overan aqueous solution of zinc acetate [146,147]. The thin ZnSe film (ca. 50–100 nmthick) formed at the gas–solution interface could be lifted up and placed on anydesired substrate. The films deposited at 80°C were found to be zincblende ZnSeby both XRD and ED (no XRD pattern was observed for films deposited at 2°C).The optical bandgap was 2.62 eV, and resistivity was 107 �-cm.

4.6 HgS

HgS possesses a very low value of Ksp (6 � 10�53) and therefore is expected to berelatively simple to deposit. In fact, apart from some relatively early literature onternary mercury sulphides with lead [148,149] and cadmium [150], which will bediscussed in Chapter 8, only three separate studies on CD HgS have been found.

Perakh and Ginsburg [151] deposited HgS films using two different tech-niques. One was a standard CD method using thiourea and HgCl2 complexed withiodide (iodide is a strong complexant for Hg2�) in an alkaline solution. The othertechnique was simply reaction of a low concentration (ca. 2 mM) of HgCl2 with(at least three times the Hg concentration) Na2S solution, which precipitates HgSas a colloid. HgS deposited slowly over many hours. It is interesting that whilesome film deposition is expected by this second method, thicknesses of at least 0.7�m were obtained—much thicker than would intuitively be expected. The tem-perature dependence of the growth (at least for the thiourea method) was compli-cated and depended on other parameters. Optimum temperatures were aroundroom temperature; temperatures higher than 25°C resulted in rough films. The ab-sorption spectra of both types of films were rather strange—a gradual absorptiononset at somewhat less than 700 nm (ca. 1.8 eV) and a sharp onset at 400 nm (3.1eV). HgS occurs in (at least) two forms: red (distorted rocksalt, cinnabar—the sta-ble form at normal temperatures) with a bandgap of ca. 2.0 eV and a black form


(zincblende) with a bandgap of ca. 0.5 eV (although large variations of this valuehave been reported—see Hg ternaries in Chap. 8). On precipitation of a solutionof Hg2� ions with sulphide solution, a black precipitate is formed (it is not clearif this is the same as the zincblende form), which eventually should convert to themore stable red form, although this change might take a long time. From the ab-sorption spectra of the films given in this study, with no absorption at wavelengthsgreater than 700 nm, the films were clearly not of the black form. The sharp onsetat 3.1 eV could conceivably be due to absorption by very small crystals However,it is more likely that it is a higher transition of HgS, since HgS that is not quan-tized also shows this feature (see later). The gradual onset from ca. 1.8 nm couldbe due to absorption in a distribution of larger sizes also containing a low con-centration of “black” heavily quantized HgS (low concentration since the absorp-tion below 2 eV is very weak). It must be stated that no structural characterizationof these films was made; however, it is unlikely that anything other than HgS wasformed under the deposition conditions.

The other studies used thiosulphate as a sulphur source. One was carried outin a simple mixture of HgCl2 and Na2S2O3, presumably under acidic conditions[152]. Crystal size measured by XRD was reported to range from 3 nm to 8 nmover a temperature range of 0–85°C (film thickness also varied over this temper-ature range from 50 to 180 nm), and this size regime was confirmed by TEM. Thebandgap estimated from optical spectra varied monotonically as a function of de-position temperature from 2.3–2.4 eV at 0°C to ca. 1.9 eV at 85°C, due to sizequantization. The resistivity of the films varied from 104 to 103 �-cm over thesame temperature range and generally decreased strongly with increasing mea-surement temperature to low values (very approximately 100 �-cm) at tempera-tures over 150°C.

A similar deposition, only carried out under alkaline conditions (with addedammonia) at pH 11 and at 65°C, was described [153]. According to this study,simple mixing of a mercuric salt solution with thiosulphate results in immediateformation of a black precipitate of HgS. By first treating the Hg2� solution withaqueous ammonia, a white precipitate formed by the following reaction:

Hg(NO3)2 � 2NH3 → (NH2Hg)NO3 � NH4NO3 (4.9)

This precipitate dissolves in thiosulphate to form a thiosulphate complex, which,in common with other metal thiosulphate complexes, decomposes when heated tothe metal sulphide (see Sec. 3.3.3). Besides direct decomposition of the thiosul-phate complex, another possibility suggested in this study is formation of sulphideion by alkaline hydrolysis of thiosulphate [Eqs. (3.20) to (3.24)] and reaction withHg2� to form HgS. The substrates were glass precoated with a very thin film ofCD PbS (presumably this improved adherence and/or homogeneity).

The films deposited by this process were golden yellow if thin (ca. 100 nm)and became red (the normal cinnabar color) if thick (ca. 500 nm). The terminal


thickness of the films was 80–90 nm; thicker films were obtained by repeated depositions.

X-ray diffraction showed that the deposit (both as a film and the precipitatedmaterial) was cinnabar HgS. The diffraction peaks were sharp (crystal size at least20 nm, possibly much greater).

Optical transmission spectra showed a sharp absorption at ca. 400 nm (3.1eV) and a weaker one at �600 nm (2.0–2.1 eV), the latter probably correspond-ing to the HgS bandgap, since little or no size quantization should be observed inthese films due to the relatively large crystal size. However, there was a gradualloss in transmission to beyond 800 nm, which was more pronounced for thickerfilms. It is not clear to what extent this is due to scattering (the thicker films werereported to be less reflecting than the thinner ones, implying increased scattering)or to absorption.

The resistivity of the films was greater than 104 �-cm (calculated from a re-ported sheet resistivity of �109 � and an assumed thickness of 100 nm).

4.7 HgSe

Three studies in total have been made on CD HgSe. In the earliest, selenosemi-carbazide was used as a Se source and the Hg was complexed with iodide [99].The deposition was carried out at 20°C. No details of the films were given otherthan that they were specular.

In the other two studies, selenosulphate was used. In one, a formamide com-plex of Hg, made by dissolving HgO in formamide, was used [154]. The solutionwas made ca. 0.5 M in NaOH, and a trace of polyvinyl pyrollidone was added. Thedeposition was carried out at room temperature. The polyvinyl pyrollidone slowedthe deposition somewhat and apparently improved film uniformity and adherenceas well as slightly increased terminal thickness (500 nm). It was noted that filmswere not obtained with the usual complexants, such as ammonia, triethanolamine,and cyanide. It is not mentioned in which way these complexants were unsuitable;ammonia and triethanolamine might be too weak, resulting in immediate precipi-tation in solution. Also, addition of ammonia to some mercuric salts tends to leadto precipitation of insoluble products. Cyanide, however, is a very strong com-plexant and would be expected to control such bulk precipitation better than for-mamide. Iodide, a strong complex for Hg2� (and successfully used to deposit HgS,as described earlier), resulted in film deposition but with poor reproducibility.

No XRD pattern was found for the films, and on this basis they were be-lieved to consist of amorphous HgSe. Based on more recent XRD studies ofnanocrystalline materials, the lack of an XRD pattern was likely due to very smallcrystal size (supported by the increased bandgap; see later). Annealing at 200°C“crystallized” the HgSe to an extent that it was clearly identified by XRD. Opti-cal spectroscopy gave a bandgap value of 1.42 eV. Bulk HgSe is a semimetal with


a bandgap of -0.15 eV (i.e., overlapping valence and conduction bands). This sug-gests very strong size quantization (see HgSe in Chap. 10). The room-temperatureresistivity was ca. 300 �-cm.

In the other study [155], ammonia-complexed Hg(NO3)2 was mixed withthe selenosulphate solution. As for the corresponding HgS deposition, a white pre-cipitate formed on addition of ammonia to the Hg(NO3)2 [Eq. (4.9)]. This precip-itate dissolved partly in the excess ammonia used, due to formation of various am-mine complexes, and completely when the selenosulphate solution was added,due to additional formation of selenosulphate (and maybe sulphite from the excesssulphite in the selenosulphite solution) complexes. It is likely that mixed ammine-selenosulphate/sulphite complexes were formed. The deposition was carried outon polyester substrates (the transparencies used in overhead projectors) at 10°C.Deposition occurred over several hours to a terminal thickness of ca. 250 nm. Bulkprecipitation occurred in parallel with the deposition, suggesting that the clustermechanism was dominant.

X-ray diffraction of the powder precipitated in solution confirmed it to beHgSe (sphalerite). The spectrum of the film showed a strong (111) peak and vir-tually no other reflection, suggesting a high degree of texture of these films. Fromthe peak broadening, a crystal size of 7.7 nm was calculated.

From optical spectra, a bandgap of 2.5 eV was estimated (based on an indi-rect gap), and this increase from the negative bandgap of bulk HgSe (see earlier)was attributed to size quantization.

The sheet resistivity was measured to be 13 k�-cm�2. Although the filmthickness was not given, based on a thickness of 250 nm, this translates into a spe-cific resistivity of �1 �-cm. Annealing the films at low temperatures (ca. 100°C)results in a decrease in resistivity up to as much as an order of magnitude (the crys-tal size, as measured by XRD, increases only slightly via this treatment).

REFERENCES

1. D Lincot, M Froment, H Cachet. In: RC Alkire, DM Kolb, eds. Adv. Electrochem.Sci. Eng., Vol. 6. New York: Wiley-VCH, 1998, p 165.

2. I Kaur, DK Pandya, KL Chopra. J. Electrochem. Soc. 127:943, 1980.3. D Lincot, R Ortega-Borges, M Froment. Philos. Mag. B 68:185, 1993.4. T Nakanishi, K Ito. Sol. Energy Mater. Sol. Cells 35:171, 1994.5. M Rami, E Benamar, M Fahoume, F Chraibi, A Ennaoui. Solid State Sci. 1:179, 1999.6. P O’Brien, T Saeed. J. Crystal Growth 158:497, 1996.7. MTS Nair, PK Nair, RA Zingaro, EA Meyers. J. Appl. Phys. 75:1557, 1994.8. DS Boyle, P O’Brien, DJ Otway, O Robbe. J. Mater. Chem. 9:725, 1999.9. ML Breen, JT Woodward, DK Schwartz, AW Apblett. Chem. Mat. 10:710, 1998.

10. S Gorer, G Hodes. J. Phys. Chem. 98:5338, 1994.11. S Kohle, SK Kulkarni, AS Nigavekar, SK Sharma. Sol. Energy Mater. 10:47, 1984.


12. ME Özsan, DR Johnson, M Sadeghi, D Sivapathasundaram, D Lincot, B Mokili, M Froment, J Vedel, LM Peter, G Goodlet, RC Walker. In: 13th ECPV Solar En-ergy Conf. Nice, France, 1995, p 2115.

13. MA Martínez, C Guillén, J Herrero. Appl. Surf. Sci. 140:182, 1999.14. R Castro-Rodríguez, AI Oliva, V Sosa, F Caballero-Briones, JL Peña. Appl. Surf.

Sci. 161:340, 2000.15. NR Pavaskar, CA Menezes, APB Sinha J. Electrochem. Soc. 124:743, 1977.16. TL Chu, SS Chu, N Schultz, C Wang, CQ Wu. J. Electrochem. Soc. 139:2443, 1992.17. PC Rieke, SB Bentjen. Chem. Mat. 5:43, 1993.18. PN Gibson, ME Özsan, D Lincot, P Cowache, D. Summa. Thin Solid Films 361:34,

2000.19. A Mondal, TK Chaudhuri, P Pramanik. Sol. Energy Mater. 7:431, 1983.20. KL Narayanan, R Rajaraman, MC Valsakumar, KGM Nair, KP Vijayakumar. Mat.

Res. Bull. 34:1729, 1999.21. WO Milligan. J. Phys. Chem. 38:797, 1934.22. NG Dhere, DL Waterhouse, KB Sundaram, O Melendez, NR Parikh, B Patnaik.

J. Mater. Sci. 6:52, 1995.23. PJ Sebastian. Appl. Phys. Lett. 62:2956, 1993.24. PJ Sebastian, M Ocampo. J. Appl. Phys. 77:4548, 1995.25. D Petre, I Pintilie, E Pentia, T Botila. Mater. Sci. Eng. B. 58:238, 1999.26. M Ichimura, F Goto, E Arai. J. Appl. Phys. 85:7411, 1999.27. A Arias-Carbajal Reádigos, VM García, O Gomezdaza, J Campos, MTS Nair,

PK Nair. Semicond. Sci. Technol. 15:1022, 2000.28. MTS Nair, PK Nair, J Campos. Thin Solid Films 161:21, 1988.29. O Vigil, Y Rodriguez, O Zelaya-Angel, C Vazquez-Lopez, A Morales-Acevedo,

JG Vazquez-Luna. Thin Solid Films 322:329, 1998.30. LP Deshmukh, SG Holikatti, BM More. Mater. Chem. Phys. 39:278, 1995.31. JM Gracia-Jiménez, G Martínez-Montes, R Silva-González. J. Electrochem. Soc.

139:2048, 1992.32. KK Nanda, SN Sarangi, S Mohanty, SN Sahu. Thin Solid Films 322:21, 1998.33. R Lozada-Morales, O Zelaya-Angel. Thin Solid Films 281–282:386, 1996.34. PK Nair, MTS Nair, J Campos, LE Sansores. Solar Cells 22:211, 1987.35. JH Schön, O Schenker, B Batlogg. Thin Solid Films 385:271, 2001.36. PK Nair, J Campos, MTS Nair. Semicond. Sci. Technol. 3:134, 1988.37. K Ito, K Shiraishi. Sol. Energy Mater. Sol. Cells 35:179, 1994.38. WJ Danaher, LE Lyons, GC Morris. Sol. Energy Mater. 12:137, 1985.39. C Guillén, MA Martínez, J Herrero. Thin Solid Films 335:37, 1998.40. M Nagao, S Watanabe. Jpn. J. Appl. Phys. 7:684, 1968.41. IO Oladeji, L Chow. J. Electrochem. Soc. 144:2342, 1997.42. O Vigil, O Zelaya-Angel, Y Rodriguez, A Morales-Acevedo. Phys. Status Solidi (a)

167:143, 1998.43. JA Akintunde. Phys. Status Solidi (a) 179:363, 2000.44. PK Nair, MTS Nair, VM Garcia, OL Arenas, Y Pena, A Castillo, IT Ayala, O

Gomezdaza, A Sanchez, J Campos, H Hu, R Suarez, ME Rincon. Sol. Energy Mater.Sol. Cells 52:313, 1998.


45. PK Nair, MTS Nair, J Campos. In: CM Lampert, ed. Proc. SPIE, San Diego, CA,1987, Vol. 823, p 256.

46. PK Nair, MTS Nair, J Campos. Sol. Energy Mater. 15:441, 1987.47. IO Oladeji, L Chow, JR Liu, WK Chu, ANP Bustamante, C Fredricksen,

AF Schulte. Thin Solid Films 359:154, 2000.48. L Pintilie, E Pentia, I Pintilie, D Petre. Mater. Sci. Eng. B 44:403, 1997.49. M Froment, MC Bernard, R Cortes, B Mokili, D Lincot. J. Electrochem. Soc.

142:2642, 1995.50. AG Shikalgar, SH Pawar. Sol. State Commun. 32:361, 1979.51. AG Shikalgar, SH Pawar. Thin Solid Films 61:313, 1979.52. CD Lokhande, SH Pawar. Sol. State Commun. 44:1137, 1982.53. JA Akintunde. J. Mater. Sci.-Mater. Electron. 11:503, 2000.54. AI Oliva, R Castro-Rodriguez, O Ceh, P Bartolo-Perez, F Caballero-Briones,

V Sosa. Appl. Surf. Sci. 148:42, 1999.55. D Bhattacharyya, MJ Carter. Thin Solid Films 288:176, 1996.56. D Lincot, R Ortega-Borges, M Froment. Appl. Phys. Lett. 64:569, 1994.57. R Cortes, M Froment, B Mokili, D Lincot. Philos. Mag. Letts. 73:209, 1996.58. D Lincot, B Mokili, R Cortes, M Froment. Microsc. Microanal. Microstruct. 7:217,

1996.59. MJ Furlong, D Lincot, M Froment, R Cortés, AN Tiwari, M Krejci, H Zogg. In: 14th

ECPV Solar Energy Conf. Barcelona, Spain, 1997, p 1291.60. MJ Furlong, M Froment, MC Bernard, R Cortés, AN Tiwari, M Krejci, H Zogg,

D Lincot. J. Crystal Growth 193:114, 1998.61. YK Hwang, SY Woo, JH Lee, DY Jung, YU Kwon. Chem. Mat. 12:2059, 2000.62. RL Call, NK Jaber, K Seshan, JR Whyte Jr. Sol. Energy Mater. 2:373, 1980.63. GA Kitaev, SG Mokrushin, AA Uritskaya. Kolloid Zh. 27:51, 1965.64. R Ortega-Borges, D Lincot. J. Electrochem. Soc. 140:3464, 1993.65. V Polpescu, EM Pica, I Pop, R Grecu. Thin Solid Films 349:67, 1999.66. K Yamaguchi, T Yoshida, T Sugiura, H Minoura. J. Phys. Chem. B 102:9677, 1998.67. K Yamaguchi, T Yoshida, N Yasufuku, T Sugiura, H Minoura. Electrochem.

67:1168, 1999.68. PK Nair, MTS Nair. Sol. Energy Mater. 15:441, 1987.69. V Weiss, AA Friesem, A Peled. Thin Solid Films 218:193, 1992.70. F Goto, M Ichimura, E Arai. Jpn. J. Appl. Phys. 2 36:L1146, 1997.71. M Ichimura, F Goto, E Arai. J. Electrochem. Soc. 146:1028, 1999.72. JGea Vázquez-Luna. J. Crystal Growth 187:380, 1998.73. M Gros-Jean, R Herino, D Lincot. J. Electrochem. Soc. 145:2448, 1998.74. DS Boyle, A Bayer, MR Heinrich, O Robbe, P O’Brien. Thin Solid Films 361:150,

2000.75. D Hariskos, M Powalla, N Chevaldonnet, D Lincot, A Schindler, B Dimmler. Thin

Solid Films 387:179, 2001.76. A Kylner, J Lindgren, L Stolt. J. Electrochem. Soc. 143:2662, 1996.77. A Kylner. J. Electrochem. Soc. 146:1816, 1999.78. A Kylner, M Wirde. Jpn. J. Appl. Phys. 1 36:2167, 1997.79. DW Niles, G Herdt, M AlJassim. J. Appl. Phys. 81:1978, 1997.


80. Y Hashimoto, T Nakanishi, T Andoh, K Ito. Jpn. J. Appl. Phys. 2 34:L382, 1995.81. T Nakada, H Fukuda, A Kunioka, S Niki. In: 13th ECPV Solar Energy Conf. Nice,

France, 1995, p 1597.82. R Jayakrishnan, JP Nair, BA Kuruvilla, SK Kulkarni, RK Pandey. J. Mater. Sci.

7:193, 1996.83. GA Kitaev, TS Terekhova. Zh. Neorg. Khim. 15:48, 1970.84. RC Kainthla, DK Pandya, KL Chopra. J. Electrochem. Soc. 127:277, 1980.85. RA Boudreau, RD Rauh. J. Electrochem. Soc. 130:513, 1983.86. GA Kitaev, AA Uritskaya, LE Yatlova, TS Terekhova, TI Dzyuba. Zhurn. Neorg.

Khim. 35:3072, 1990.87. G Hodes, A Albu-Yaron, F Decker, P Motisuke. Phys. Rev. B 36:4215, 1987.88. S Gorer, G Hodes, Y Sorek, R Reisfeld. Mater. Lett. 31:209, 1997.89. F Trojanek, R Cingolani, D Cannoletta, D Mikes, P Nemec, E Uhlirova, J Rohovec,

P Maly. J. Cryst. Growth 209:695, 2000.90. G Hodes. Isr. J. Chem. 33:95, 1993.91. R Garuthara, G Levine. J. Appl. Phys. 80:401, 1996.92. P Nemec, D Mikes, J Rohovec, E Uhlirova, F Trojanek, P Maly. Mater. Sci. Eng.

B 69:500, 2000.93. M Perakh, A Peled. Thin Solid Films 50:293, 1978.94. H Cachet, H Essaaidi, M Froment, G Maurin. J. Electroanal. Chem. 396:175, 1995.95. KC Mandal, O Savadogo. J. Mater. Sci. Lett. 10:1446, 1991.96. R Lozada-Morales, M Rubin-Falfan, O Portillo-Moreno, et al. J. Electrochem. Soc.

146:2546, 1999.97. MTS Nair, PK Nair, H Pathirana, RA Zingaro, EA Meyers. J. Electrochem. Soc.

140:2987, 1993.98. BK Rai, HD Bist, RS Katiyar, MTS Nair, PK Nair, A Mannivannan. J. Appl. Phys.

82:1310, 1997.99. AA Velykanov, EK Ostrovskaya, NP Garina, VA Turacova, AA Tchurkan. Ukr.

Chim. Zh. 49:764, 1983.100. H Cachet, R Cortés, M Froment, G Maurin, N Shramchenko. J. Electrochem. Soc.

144:3583, 1997.101. E Lifshitz, I Dag, I Litvin, G Hodes, S Gorer, R Reisfeld, M Zelner, H Minti. Chem.

Phys. Lett. 188, 1998.102. XC Ai, R Jin, CB Ge, JJ Wang, YH Zou, XW Zhou, XR Xiao. J. Chem. Phys.

106:3387, 1997.103. P Maly, J Kudrna, F Trojanek, D Mikes, P Nemec, AC Maciel, JF Ryan. Appl. Phys.

Lett. 77:2352, 2000.104. K Rajeshwar, L Thompson, P Singh, RC Kainthla, KL Chopra. J. Electrochem. Soc.

128:1744, 1981.105. RC Kainthla, DK Pandya, KL Chopra. Sol. State Electron. 25:73, 1982.106. VM Garcia, MTS Nair, PK Nair, RA Zingaro. Semicond. Sci. Technol. 11:427, 1996.107. MTS Nair, PK Nair, RA Zingaro, EA Meyers. J. Appl. Phys. 74:1879, 1993.108. G Hodes, IDJ Howell, LM Peter. In: Photochemical and Photoelectrochemical Con-

version and Storage of Solar Energy. ZW Tian, Y Cao, eds. Int. Acad. Publishers,Beijing, China. 1993, p 331.


109. L Kronik, L Burstein, M Leibovitch, Y Shapira, D Gal, E Moons, J Beier, G Hodes,D Cahen, D Hariskos, R Klenk, HW Schock. Appl. Phys. Lett. 67:1405, 1995.

110. E Lifshitz, I Dag, I Litvin, G Hodes. J. Phys. Chem. B 102:9245, 1998.111. B Alperson, I Rubinstein, G Hodes. Phys. Rev. B 6308:1303, 2001.112. GK Padam, SK Gupta. Appl. Phys. Lett. 53:865, 1988.113. RW Buckley. In: 11th ECPV Solar Energy Conf. Montreux, Switzerland, 1992, p

962.114. VB Patil, PD More, DS Sutrave, GS Shahane, RN Mulik, LP Deshmukh. Mater.

Chem. Phys. 65:282, 2000.115. AV Kalyakina, RI Pelyukpashidi. Tr. Khim. Met. Inst. Akad. Nauk. Kaz. SSR

17:114, 1973.116. M Sotelo-Lerma, RA Zingaro, SJ Castillo. J. Organomet. Chem. 623:81, 2001.117. P O’Brien, J McAleese. J. Mater. Chem. 8:2309, 1998.118. B Mokili, M Froment, D Lincot. J. de Phys. IV 5:261, 1995.119. P O’Brien, DJ Otway, D Smith-Boyle. Thin Solid Films 361:17, 2000.120. JM Doña, J Herrero. J. Electrochem. Soc. 141:205, 1994.121. IO Oladeji, L Chow. Thin Solid Films 339:148, 1999.122. B Mokili, Y Charreire, R Cortes, D Lincot. Thin Solid Films 288:21, 1996.123. J McAleese, P O’Brien. In Mater. Res. Soc. Symp. Proc., 1998; Vol. 485; p 255.124. IC Ndukwe. Sol. Energy Mater. Sol. Cells 40:123, 1996.125. S Biswas, P Pramanik, PK Basu. Mater. Lett. 4:81, 1986.126. R Ortega Borges, D Lincot, J Vedel. In: 11th ECPV Solar Energy Conf., Montreux,

Switzerland, 1992, p 862.127. PK Nair, MTS Nair. Semicond. Sci. Technol. 7:239, 1992.128. PK Nair, MTS Nair, O Gomezdaza, RA Zingaro. J. Electrochem. Soc. 140:1085, 1993.129. OL Arenas, MTS Nair, PK Nair. Semicond. Sci. Technol. 12:1323, 1997.130. K Yamaguchi, T Yoshida, T Sugiura, H Minoura. J. Mater. Res. 13:917, 1998.131. CD Lokhande. Mater. Chem. Phys. 28:145, 1991.132. M Ichimura, F Goto, Y Ono, E Arai. J. Cryst. Growth 199:308, 1999.133. GA Kitaev, AA Uritskaya, LE Yatlova, VR Mirolyubov. Russ. J. Appl. Chem.

67:1415, 1994.134. QW Chen, YT Qian, ZY Chen, L Shi, XG Li, GE Zhou, YH Zhang. Thin Solid Films

272:1, 1996.135. QW Chen, XG Li, Y Qian, JS Zhu, G Zhou, WP Zhang, YH Zhang. Appl. Phys.

Lett. 68:3582, 1996.136. P Pramanik, S Biswas. J. Electrochem. Soc. 133:350, 1986.137. RK Nkum, AA Adimado, H Totoe. Mater. Sci. Eng. B 55:102, 1998.138. GA Kitaev, TP Sokolova. Russ. J. Inorg. Chem. 15:167, 1970.138a. GA Kitaev, TP Sokolova. USSR Patent 356, 319 1972.139. CD Lokhande, PS Patil, A Ennaoui, H Tributsch. Appl. Surf. Sci. 123:294, 1998.140. JM Doña, J Herrero. J. Electrochem. Soc. 142:764, 1995.141. AM Chaparro, MA Martinez, C Guillen, R Bayon, MT Gutierrez, J Herrero. Thin

Solid Films 361:177, 2000.142. A Ennaoui, M Weber, M Saad, W Harneit, MC Lux-Steiner, F Karg. Thin Solid

Films 361:450, 2000.


143. J Herrero, MT Gutierrez, C Guillen, JM Dona, MA Martinez, AM Chaparro, R Bayon. Thin Solid Films 361:28, 2000.

144. CA Estrada, PK Nair, MTS Nair, RA Zingaro, EA Meyers. J. Electrochem. Soc.141:802, 1994.

145. C Wang, XF Qian, WX Zhang, XM Zhang, Y Xie, YT Qian. Mat. Res. Bull.34:1637, 1999.

146. GN Chaudhari, SD Sathaye, P Singh, VJ Rao, V Manorama, RS Bhide, SV Bho-raskar. J. Mater. Sci. Lett. 11:1097, 1992.

147. GN Chaudhari, S Manorama, VJ Rao. J. Phys. D:Appl. Phys. 25:862, 1992.148. NC Sharma, DK Pandya, HK Sehgal, KL Chopra. Mat. Res. Bull. 11:1109, 1976.149. NC Sharma, DK Pandya, HK Sehgal, KL Chopra. Thin Solid Films 42:383, 1977.150. M Skyllas-Kazacos, JF McCann, R Arruzza. Appl. Surf. Sci. 22/23:1091, 1985.151. M Perakh, H Ginsburg. Thin Solid Films 52:195, 1978.152. SS Kale, CD Lokhande. Mater. Chem. Phys. 59:242, 1999.153. M Najdoski, I Grozdanov, SK Dey, B Siracevska. J. Mater. Chem. 8:2213, 1998.154. P Pramanik, S Bhattacharya. Mat. Res. Bull. 24:945, 1989.155. B Pejova, M Najdoski, I Grozdanov, SK Dey. J. Mater. Chem. 11:2889, 1999.


5PbS and PbSe

5.1 HISTORICAL BACKGROUND

PbS holds the honor of being the first reported compound to be deposited by CD.In 1869, Puscher described a “new and cheap process, without using dyes, to coatvarious metals with splendid lustrous colors” [1]. This involved deposition froma thiosulphate solution of lead acetate (and also, in the same paper, from Cu andSb salts to give presumably corresponding sulphides). These shiny, colored coat-ings prompted further studies in this process, both to expand the process to othermetal sulphides and to understand the process. These studies are discussed in Sec-tion 5.2.1.

The common thiourea process for CD was also first used for PbS [2]. Thethiourea method became the preferred one for the emerging development of PbSphotoconductive cells as infrared (IR) detectors during the Second World War.Obviously much of this early work, carried out by German groups for military pur-poses, was secretive and was not published at the time. Later, photoconductivecells using CD PbS (which gave better cells than did the more conventional, evap-orated PbS films [3]) became commercial as IR detectors, and, together with CDPbSe, they remain so until this day. An early description of such cells is given byKicinski [4]. This application then became the driving force for CD studies, whichwere limited almost entirely to PbS and PbSe up to the start of the 1960s. Section5.2.6 gives a brief overview of the operation of photoconductive cells in generaland lead salt cells in particular.


Since the expansion of CD to II–VI compounds in the 1960s and subse-quently to other compounds, work on the IV–VI materials has not kept pace withthis expansion, and there has been only a relatively limited amount of work sincethen, mainly on the use of PbS for solar control coatings. This chapter does not as-pire to comprehensively cover the older literature, although it is intended that theimportant aspects of these studies be included. Previous reviews that deal withthese materials (mostly reviews on infrared detectors) include Refs. 3 and 5–8.

5.2 PbS

5.2.1 Deposition Using ThiosulphateThe original CD process involves decomposition of a thiosulphate solution ofPbAc2 [1]. From the outset, it must be noted that the mechanism of this depositionhas not been unambiguously elucidated up to the present time. There are two mainpossibilities for the reaction.

1. Decomposition of a thiosulphate complex of Pb, a possible reaction being

Pb(S2O3)22� � H2O → PbS � S2O3

2� � H2SO4 (5.1)

2. Decomposition of S2O32� to give free sulphide ion, which then reacts

with free Pb2�. It has been suggested [9,10] that this can occur by reduction of el-emental S, which forms in acidic solutions of S2O3

2�

S2O32� � H� → S � HSO3

� (5.2)

by electrons formed in the half cell reaction:

2S2O32� → S4O6

2� � 2e� (5.3)

(An internal electrochemical mechanism was proposed long ago for deposition oncertain metal substrates, since the rate of deposition sometimes depended on thenature of the substrate [11].) The standard potential of Reaction (5.3) is �0.08 V,considerably more positive than the reduction potential of S to S2� (�0.45 V).Free sulphide, if formed, would be in a very low concentration, since it will be re-moved continually by precipitation of PbS; this will move the S reduction poten-tial strongly positive according to the Nernst equation [Eq. (1.32)]. This positiveshift will be even greater than normal because of the non-Nernstian behavior ofthe S2�/S couple when [S] � [S2�] (at least in alkaline solution) [12]. In opposi-tion to this, the solubility of S in the (slightly acidic) aqueous solutions is very low,which will move the potential in the opposite direction. Add to this the very smallconcentration of S2� in acid solution [Eq. (1.15)], and it becomes clear that it isnot trivial to estimate the feasibility of the formation of PbS by free sulphide. Thenon-Nernstian behavior of the sulphur-rich S2�/S couple and the lack of knowl-edge of the solubility of free S in the deposition solution are the two factors thatcomplicate what would have been a tractable thermodynamic calculation.


The decomposition of metal thiosulphates to the corresponding sulphide iswell known and is generally assumed to occur by breaking of the S–S bond in thio-sulphate. Thus, it has been shown, by a radioactive tracer method, that if thiosul-phate is prepared by dissolving radioactive S in sulphite solution:

S* � SO32�D S*MSO3

2� (5.4)

and this solution is reacted with Ag� to form a precipitate of Ag2S2O3, then theAgS formed on boiling this precipitate in water contains only labeled S, showingthat the metal bonds only to the S which is not attached to oxygen [13]. Admit-tedly, even if the AgS was formed by free sulphide, the same result would be prob-able due to the lability of the SMS bond; however, such reactions have been stud-ied for many metal cations, and it is invariably assumed (and this does notnecessarily make the assumption correct) that the reaction occurs by simple fis-sion of the SMS bond in the thiosulphate. Additionally, since this reaction occursmost readily for metals with a high affinity for S (Ag, Cu(I), Hg and Pb), there isno reason to expect that the MMS bond present in the thiosulphate (or complexthiosulphate) would be broken, as would be the case were the formation of metalsulphide to occur by free sulphide, although, as has already been pointed out, thesesame metal sulphides have a very low-solubility product and therefore very littlesulphide is required to precipitate them.

One potentially useful piece of information that can be explained more read-ily based on a free-sulphide generation comes from an early study on the forma-tion of PbS by boiling Pb2� and thiosulphate in water, when it was found that PbSformed much more readily when excess thiosulphate was present [14].

To sum up, the mechanism of formation of PbS using thiosulphate is still notdefined, either in general or even for any specific case. Some dedicated researchto solve this question is clearly required.

While a number of early studies described PbS (and other sulphides) for-mation from thiosulphate solutions [1,11,15,16], these early studies provided lit-tle characterization of the films other than the interference colors obtained (due todifferent thicknesses of the films). There appear to be only two modern studiesthat provide some general characterization of these PbS films. In one [10], filmswere deposited on glass at 80°C from a solution of PbAc2 and Na2S2O3 at a pHbetween 5 and 6. X-ray diffraction (XRD) showed sharp peaks of polycrystallinePbS. Although the thickness of the films was not given, they were clearly thick byCD standards, since the optical transmission over the range from 400 to 2000 nmwas less than 1%. Electrical resistivity was measured to be ca. 105–106 �-cm. Inanother study [17] (also by the same group on Cu2S-coated glass [18]), the condi-tions were similar (the pH range was between 4 and 6, and acetic acid was added,if necessary, to reach this pH range; the bath temperature was slightly lower—ca.70°C—and the reactant concentrations were lower). The PbS films were thinner(0.1 �m), as could be seen by the greater optical transmission (�60% at 2.5 �m,


dropping gradually to ca. 7% at 500 nm and more rapidly at shorter wavelengths).Such semitransparent films are considered useful for solar control coatings (seeSec. 2.13). The films were p-type with a very low resistance (4 k�/sq, equivalentto 4 � 10�2 �-cm). The large difference between the resistivities in the two stud-ies seems worthy of further investigation.

5.2.2 Deposition Using Thiourea

With the exception of the few cases just described, all the modern (and most of theearlier) studies on PbS deposition were carried out using thiourea as a source ofsulphide, as first described by Emerson-Reynolds in 1884 [2]. In the originalstudy, the lead was present as a strongly alkaline solution of lead tartrate (proba-bly a mixed tartrate–hydroxide complex of lead). When this solution was heatedwith a solution of thiourea, “at about 50° a specular layer forms at the bottom andsides of the vessel. When the beaker is thoroughly clean in the first instance, theadhesion is uniform and strong.” Most of the PbS formed as a precipitate. Thecompeting effects of temperature were noted: At room temperature the depositionwas very slow (days) and the film tended to be less even (more heterogeneous tothe eye), while at higher temperatures thin (due to excessive precipitation) brown-ish films were obtained. Deposition was reported on a number of different mate-rials (including porcelain and ebonite as well as iron and steel, although films onthe metals had poorer adhesion).

The same basic technique was used to study the deposition of PbS on dif-ferent types of glass substrates. The quality of the film varied greatly from verypoor, partially formed films that were not adherent (quartz and borosilicate glass),through occasionally good but irreproducible films (on sodalime glass), to homo-geneous, adherent films on flint (lead containing) and Zn-containing (crown)glasses [19]. The good adherence on the Pb and Zn glasses was explained by for-mation of insoluble sulphides by these metals; such sulphides would form a goodbinding site for further deposition of PbS (presumably also for other compounds).

Lead acetate was later usually employed as the Pb salt (e.g., Refs. 4 and 20).The thiourea acts not only as a source of S but also as a complexing agent, asshown by Kicinski. In the absence of thiourea, hydroxide can also complex Pb2�,but a considerable excess is usually required in practice. Small amounts of leadoxide and hydroxide were detected in these films by XRD [4].

There does not seem to be a clear consensus as to the mechanism of PbS formation in these and similar studies. It is often stated that the formation of PbSoccurs via decomposition of a Pb–thiourea complex species [4,21,22]. This wasoften based on the absence of any measurable concentration of sulphide on alka-line hydrolysis of thiourea. However, as discussed in Section 3.3.3.1, this is not avalid criterion for the absence of a sulphide-mediated reaction. Even today, it can-not be stated categorically which mechanism is operative or even dominant.


A microscopic study of the PbS formed in these films deposited on glass at50°C showed the formation of large (0.2–1 �m) cubic crystals from relativelyconcentrated solutions (the films were specularly reflecting) and somewhat larger,but less well-defined crystals from more dilute solutions (resulting in more highlyscattering films) [23]. Addition of very small amounts of CuSO4 to the dilute de-position solution resulted in specularly reflecting films with smaller, more evenlysized crystals. From XRD line broadening, the crystal size of PbS deposited by themethod described by Kiciniski was estimated at ca. 70 nm [24]. Another report us-ing films prepared by Kicinski on pyrex described a preferred orientation with the(001) faces parallel to the glass surface and a crystal size larger than 15 nm and insome cases ca. 50 nm [25].

5.2.3 Variations in Deposition

This section deals only with thiourea-based baths. There is little variation in thethiosulphate baths that have been reported.

5.2.3.1 NH4OH (in place of NaOH or KOH)

Ammonium hydroxide has also been used in place of NaOH or KOH [20,24,26].In reported contrast to films deposited from alkali metal hydroxide, these films,prepared at or slightly above room temperature, were photoconductive (photosen-sitivity ca. 10) as deposited without need for air-annealing [26]. The crystal sizeof films deposited at different temperatures was measured (XRD line broadening)to be 10–15 nm (30°C), 17 nm (40°C), and 39 nm (50°C) [24,26]. The presenceof strain in the crystals was inferred from the same XRD measurements [24].

5.2.3.2 Addition of Halides

A comprehensive study has been made on the effect of added ammonium halidesto deposition from solutions of citrate-complexed Pb2� containing NH4OH [27].The effects of the ammonium halides were both strong and varied. The depositionrate decreased with increased halide concentration. This is expected, if only dueto the lower pH of the buffered solution. Also, the retardation effect of the am-monium halide increased from Cl to I. This could be due to partial removal of Pbby the sparingly soluble halides (the iodide is the least soluble and therefore willmost effectively remove the Pb). It is also possible that the more strongly adsorbediodide retards growth due to adsorption and surface capping. This is supported bythe gradual decrease in crystal size (measured by electron microscopy) from ca.0.8 �m (no halide) through 0.5 �m (Cl), 0.3 �m (Br), and 0.2 �m (I), which par-allels the increase in adsorption power of the halide ions from Cl to I. On the otherhand, substitution of NH4I by KI or NaI does not have the same effect (the crys-tal size remains ca. 0.8 �m). The films were preferentially (111) textured in theabsence of ammonium halide, and this texture decreased as the ammonium halide


concentration, as well as its atomic weight, increased, again as expected based onadsorption of halide; but again this texture-reducing effect did not occur (at leastnot to the same extent) with KI or NaI. Other effects of the different halides, givenonly for the ammonium halides (i.e., it was not known whether NaI or KI had asimilar effect or not), were change in shape of the crystals from cubic (Cl) tospherical (I) (which could also be explained by strong adsorption of iodide on allcrystal faces) and an energy of activation of the deposition varying from 38 (Cl)to 67 (I) kJ/mole; Both are characteristic of a chemical rate-determining step, butit is clear that there is an important difference in the two mechanisms. It seemsthat, while anion adsorption may play a role in these effects, the main role, at leastin those effects where ammonium and alkali metal halides were compared, has an-other explanation. A possible one, in particular for the effect on deposition rate, isthe lower pH of the ammonium-buffered solutions.

Another property of the PbS films deposited from the ammonium-bufferedsolutions was the relatively high photosensitivity obtained for the as-depositedfilms, particularly from the Cl bath. Increases in photosensitivity of close to two orders of magnitude were obtained, compared to films deposited from solutionscontaining no ammonium halide. It was suggested that the halide compensatedthe PbS (low-temperature measurements showed quasi-intrinsic conduction inthese films), resulting in high photosensitivity. Additionally, there is a generalinverse relation between the photosensitivity and the grain size/degree of textur-ing. This relationship between grain size and photosensitivity is often seen andimplies a major role of the grain surfaces in the photoconduction mechanism.Note that this is in contrast to photoconductivity in the II–VI semiconductors,where there was no obvious correlation between sensitivity and grain size (seeChap. 4).

A related study, using KBr as an additive [28], found a modest increase incrystal size with increasing KBr concentration together with an increase in crys-tal height (up to 0.4 mM Br concentration followed by a subsequent decrease withfurther increase in KBr concentration) and a slight (200) preferred orientation. Itwas suggested that this crystal growth was due to retarded growth of small nucleidue to the complexing power of the bromide. Note that the pH of this bath was ca.12, probably higher than the buffered ammonium bath described earlier.

As with the ammonium halide baths, the photosensitivity of the resultingPbS films was found to increase with added KBr in the bath (no Br was foundin the layers themselves, although very small “doping” concentrations might notbe detected in the analyses), up to a maximum of 0.4 mM KBr, and then de-creased, which correlates with the crystal height. The increase in photosensitiv-ity was attributed to disorder at the grain boundaries. The study using ammo-nium halides also attributed a major role of the grain surface to thephotoconductivity, but the relation between grain size and sensitivity is very dif-ferent in the two studies.


5.2.3.3 Oxidizing Agents in the Bath

PbS films deposited from most basic baths are only weakly photoconducting. Formaximizing photoconductivity response, air-annealing is usually carried out. Byemploying an oxidizing agent in the bath, films can be deposited that have rela-tively high response as deposited (explained by introduction of sensitizing centersin the PbS; see Sec. 5.2.6). One study [29] has described the affect of an oxidizingcomponent in the bath on the sensitivity of the resulting films (ca. 1 �m thick). Un-fortunately, neither the bath conditions (probably a Pb2�/hydroxide/thiourea bath)nor the identity of the oxidizing agent was revealed, other than that they were “vari-ations of a commercially available material.” The films were all p-type with car-rier density between 1016 and 1017/cc. The detectivity of a film deposited withoutthe oxidant was lower (by orders of magnitude at room temperature, less at lowtemperatures) than for a film prepared with an excess of oxidant or a standard (pre-sumed commercial) film. The absorption spectra of the films deposited with andwithout oxidant were also quite different (see Sec. 5.2.5). It is notable that this dif-ference in absorption was not reflected in the photoconductivity response shapenear the response onset; the samples behaved similarly with a photoconductivityonset of ca. 0.4 eV, although at higher photon energies the response of the oxidant-free film decreased strongly with increasing photon energy, in contrast to a mod-est decrease for the oxidant samples. It was also noted that the film with a high ox-idant concentration appeared more porous (in SEM micrographs) than theoxidant-free film and that the particle size was ca. 70 nm (somewhat less in the ox-idant-containing films than in the oxidant-free ones). A similar study on the effectof (the as-usual unspecified) oxidant on PbS films deposited from a thiourea/hy-droxide bath found similar effects to the foregoing study, namely, better sensitiv-ity, smaller particle size (from almost a micron in the absence of oxidant droppingto �0.2 �m with a high oxidant concentration) and more porous with increasingoxidant concentration [30]. In this case, the sensitivity peaked at a certain oxidantconcentration, followed by a decrease. Also, while the oxidant films were p-type,those deposited without the oxidant were n-type. The oxygen content was constantand independent of the oxidant concentration (as found also in the previous study[29]), and it was concluded that the origin of the active sensitizing centers was re-lated not to the oxidizing agent itself but to the structure of the film.

Apart from considerations of photoconductivity, one oxidant, H2O2, wasshown to exert considerable influence on the crystallographic texture of PbS de-posited from a PbAc2/NaOH solution with thiourea at close to room temperature[31]. While the (111) reflection was fairly dominant in the absence of H2O2, the(200) reflection became very dominant in its presence. The H2O2 was added, notwith the other reactants, but some time after film formation became visible, andthe degree of texturing reached a maximum at a certain time, after which it againdecreased. No explanation for these effects was given.


5.2.3.4 Miscellaneous Variations in Deposition

Addition of Inhibitors. Addition of salts of Sb, Sn, or As to a basicPbAc2/thiourea/OH bath resulted in a slower deposition, typically by 2–4 times[21]. These metal ions all form complex sulphides in aqueous solution, and it wasbelieved that they caused “peptization” of growing nuclei; peptization convertslarge particles into colloidal ones. Probably what is meant is some sort of cappingaction due to adsorption of the retarding ions on the growing PbS crystals, pre-venting (more correctly, retarding) further growth.

Illumination-Induced Growth. As for CdS, illumination has been found toincrease the growth rate of CD PbS films [32]. The rate increase was attributed toa combination of heating by light absorption and the formation of charge carriersin the PbS film, the latter resulting in activation of the deposition. As for CD ofCdSe under illumination, described in Chapter 4, such activation may result fromeither a photochemical reaction by the photogenerated electrons and/or holes onthe PbS surface, as suggested earlier (the most likely scenario), and/or by a changein the electrical double layer at the surface of the PbS particles (adsorbed on thesurface or as a colloid in the solution near the substrate), which might lower a po-tential barrier to adsorption and coalescence of the PbS colloids.

Alkali-Metal-Free Solutions. Films of CD PbS are usually p-type as de-posited. One early suggestion to explain this was that the alkali metal ions used in the deposition solution (as NaOH or KOH) act as a p-type dopant [33]. Basedon this supposition, Bloem deposited PbS from a solution of PbAc2, hydrazine hydrate, and thiourea (free of Na or K). The as-deposited films were initially n-type but changed to p-type on exposure to air. Attempts to dope the films perma-nently n-type by adding trivalent ions to the deposition solution were unsuccess-ful. However, by depositing the films on a substrate coated with trivalent ions(such as Al, In, Ga), n-type behavior could be maintained for a considerable time.PbS p-n junctions were fabricated using this approach (see Chap. 9).

5.2.4 Substrate Effects (See also Sec. 5.2.2)

5.2.4.1 Epitaxial Deposition

There have been a few reports on epitaxial deposition of PbS on various single-crystal substrates. PbS (n-type) was epitaxially deposited on (111) Ge (5.4% mis-match) from a Pb(NO3)2/KOH/thiourea solution at room temperature with (111)orientation [34] (although another study using apparently the same conditionsfound the deposit to be p-type and polycrystalline with some (100) preferred ori-entation [35]). From a similar solution (with addition of some ethanol), PbS wasdeposited on single-crystal CdS (ca. 6.6% mismatch) with varying degrees of epi-taxy [36]. On the (0001) faces of CdS, the growth was (111) [(111) cubic corre-


sponds to (0001) hexagonal] and on the (112�0) face, it was (220) oriented. The de-gree of epitaxy was moderate on the Cd (0001) face, better on the S (0001�) face,and very high on the (112�0) face.

Using the same reactants (but at higher concentrations), also at room tem-perature, epitaxial PbS growth was observed on (100) InP, (100) Ge, and also (although with rougher morphology) on (111) Ge [37]. In these experiments, thefilms, ca. 300 nm thick, were all p-type. The energy of activation for film forma-tion in this study was 65 kJ/mole—similar to CdS deposition from thiourea solu-tions but different from the value of activation energy measured for ZnS deposi-tion from a thiourea bath (see Sec. 4.4.1). Homogeneous precipitation alsooccurred gradually, but the films were removed before it interfered by adhesion ofclusters (presumably before it became excessive).

PbS films grown on single-crystal GaAs (lattice mismatch ca. 5%) werepolycrystalline or only somewhat oriented [38], and those on Si (9% mismatch)were polycrystalline [39]. Some of these heterojunctions with CD PbS are dis-cussed in Chapter 9.

5.2.4.2 Deposition on Monolayers

PbS has been deposited from hydroxide/thiourea solutions onto Au and thiolmonolayer–coated Au substrates [40]. The quality and texturing of the films var-ied according to the specific substrate used and also on the hydroxide concentra-tion. For dilute OH� solutions (0.1 M OH�; 0.01 M Pb2�), good deposition wasobtained only on some surfaces, while for more concentrated OH� (0.25 M OH�;0.01 M Pb2�), the deposition was not very surface dependent. Various degrees of(111) or (100) texture were obtained, depending on the substrate. On bare (111)Au, a (100) texture was obtained. It was noted that a high degree of texture wasobtained on some of the monolayers, which themselves were poorly ordered. Thecrystal size varied according to both film thickness and other deposition condi-tions. For very thin films, very small crystal sizes were obtained (from a fewnanometers up). For thick films, relatively large, more or less well-defined cubiccrystals of between 50 and several hundred nanometers were obtained. In general,the depositions from the more concentrated OH� solutions gave larger, better-defined crystals. The differences were rationalized in terms of hydroxide and ion-by-ion mechanisms predominating in the low- and high-concentration OH� solu-tions, respectively.

The ability of certain monolayer-coated surfaces to enhance or retard filmgrowth was exploited to pattern CD PbS films (see also Secs. 2.8 and 4.1.5.3).This patterning was based on the UV-induced oxidation of a thiol linkage to Au(a strong bond) to a weakly bound sulphonate group that could be rinsed away[41]. Figure 5.1 shows the processes involved in this patterning. A long-chainmercapto-carboxylic acid (16-mercaptohexadecanoic acid) was self-assembled ona Au substrate and exposed to short-wavelength UV radiation through a patterned


mask, in this case a TEM grid (a). The parts of the monolayer exposed to the UVwere oxidized to sulphonates (b) that were then rinsed away, leaving the substratepatterned with mercapto-carboxylates and exposed Au (c). A long-chain alkylthiol (16-mercaptohexadecane) was then self-assembled onto the (exposed Auportions of the) substrate (d). Chemical deposition of PbS (Pb2�, NaOH, thiourea)onto this patterned substrate resulted in PbS deposition on the hydrophilic car-boxylate endgroups and almost not at all on the hydrophobic alkyl endgroups (e).Figures 5.2a and b show the resulting deposit: A dense PbS deposit formed on theparts that were unexposed to the UV radiation (dark grid areas in 5.2a and lowerpart of 5.2b), while only scattered particles were found on the exposed areas.

5.2.4.3 Ferroelectric Substrate

Deposition of PbS onto a poled ferroelectric substrate (a complicated oxide ofmainly Pb, Zr, and Ti) from a Pb(NO3)2/NaOH/thiourea bath (containing also hy-

FIG. 5.1 Scheme for patterned deposition of PbS. (a) Self-assembled monolayer (SAM)of long-chain mercapto-carboxylic acid. On exposure to UV radiation through a mask, theexposed thiol group is oxidized to a sulphonate group (b) that is weakly bound to theAu substrate and can be easily rinsed away (c). Subsequent formation of a long-chain alkylthiol SAM occurs only on the exposed Au (d). CD of PbS occurs only on hydrophilic car-boxylate endgroups and not on hydrophobic methyl groups (e). (See Ref. 41).


droxylamine hydrochloride as a reductant and a trace of Bi3�; the reason for theseadditives was not given) resulted in films with larger particle size (1 �m) com-pared to that of films deposited onto an unpoled ferroelectric or glass (particle sizeca. �0.3 �m) [42]. The larger particle size was explained by the electric field pre-sent at the substrate surface, which attracts ions from the solution and increasesthe growth rate. It was not stated whether the film thickness was greater for thepoled substrates (as should be for a faster growth rate). In any case, a faster growthrate does not necessarily mean a larger crystal size; the opposite is often true.

The electrical properties of the films also depended on the poled state of thesubstrate. The resistivity of the PbS on the poled surfaces (10 �-cm on the posi-tively poled face and 20 �-cm on the negatively poled face) was overall lower thanthat on the unpoled ceramic (varied between 10 and 100 �-cm) or on glass (50 �-cm). The photoconductivity response varied over almost an order of magnitude, de-pending on substrate and temperature. The room-temperature response was high-est on glass and lowest on the poled substrate, but this order was reversed at liquidN2 temperatures. The films were all p-type except for that deposited on the nega-tively poled face, which was n-type. All these results were explained by the effectsof the electric field and the surface charge on the depositing film. The temperaturedependence of the photoconductivity effects was attributed to trapping dominatingthe photoconductivity at lower temperatures. It should also be noted that the parti-cle size may also affect the electrical properties; e.g., fewer grain boundaries wouldresult in a lower resistivity (assuming identical bulk properties).

5.2.4.4 Deposition on Liquid

Although not CD in the usual sense of the technique, it is worth mentioning a re-port of formation of PbS films by passing H2S over PbAc2 aqueous solutionslightly acidified with acetic acid [25]. The films were picked up on a gauze and

FIG. 5.2 SEM image of (a) PbS (dark areas) growing on patterned carboxylate-termi-nated regions of substrate and (b) boundary between PbS-covered carboxylate region(lower part of micrograph) and almost bare thiol-terminated region. (From Ref. 41 withpermission from Elsevier Science).


were analyzed by electron diffraction. The (001) faces were found to be more orless parallel to the solution surface, and the crystal size was ca. 25 nm. It is clearthat this technique should be applicable to almost all metal sulphides (also se-lenides—see Sec. 5.3—and tellurides).

Deposition from colloidal sol, a related deposition technique (although notdepending on a particular substrate surface) also resulted in PbS films usinggaseous H2S. A very thin PbS film was formed when a quartz plate was immersedovernight in a solution of Pb(ClO4)2 and poly(vinyl alcohol) through which H2Shad been bubbled [43]. The absorption spectrum of this film was similar to that ofthe PbS sol and consisted of several absorption peaks with an absorption onset ofca. 630 nm (strongly blue shifted from the PbS bulk bandgap). The XRD crystalsize of the precipitate was ca. 3 nm (see Chap. 10 for more details).

5.2.5 Optical Absorption/Transmission/Reflectionof PbS Films

Some general remarks on the optical absorption of lead chalcogenides are in orderhere, since transmission spectra of thin films of these compounds are open to mis-interpretation. If a transmission spectrum of, say, a thin (�100 nm) PbS film istaken, a sharp drop in transmission will be seen in the red–near IR region (usuallybetween 600 and 800 nm). It is easy to translate this into the bandgap of PbS, andthis has been done, even in careful studies. Thus, one thorough study of films of thePb chalcogenides deposited both by CD and by evaporation (for PbTe, only byevaporation) has reevaluated their bandgaps upwards (e.g., 1.3 eV for PbS insteadof 0.4 eV) [44]. It was later clarified that the bandgaps were indeed much lowerthan originally believed (e.g., Ref. 45). These materials have an absorption coeffi-cient in the region of their bandgap, and for considerably higher photon energies,that is only moderate (ca. 104 cm�1), and only at much higher photon energies doesthis absorption coefficient increase substantially (by about an order of magnitude),giving the apparent bandgap onset measured in thin films. In thick films (micronsand up), the absorption close to the bandgap is high. However, for much thinnerfilms, this absorption is weak. Furthermore, it is often masked by reflection.

It is particularly important when dealing with films that absorb weakly andthat have high reflectivity (such as PbS) to stress that transmission measurementsthat are uncorrected for reflection should not be directly converted to absorption,since the differences in transmission with change in wavelength may well be dueto differences in reflection rather than in absorption. Since the lead chalcogenideshave high dielectric constants (therefore high refractive indices and high reflec-tion), masking of weak or even moderate absorption by reflection is probable.

Although many studies of the optical properties of CD PbS films have beenmade, most of them do not extend to wavelengths corresponding to the bulk PbSbandgap (ca. 3 �m). The study noted by Gibson [44], which was corrected for re-


flection, showed a strong absorption beginning at ca. 800 nm and a relatively weakabsorption extending out to at least 6 �m.

In another study that does show the extended IR region for 1-�m-thickfilms, the absorption edge is very dependent on the presence or absence of an ox-idizing agent in the bath (the absorption coefficients at energies well above the on-set were similar) [29]. The films prepared with oxidant had a very gradual onsetstarting well below the normal bandgap of PbS (ca. 0.4 eV), while that depositedwithout oxidant exhibited a sharp onset at ca. 0.52 eV, considerably above the nor-mal bandgap. Various possibilities might explain these differences. The lower on-set of the “oxidant” films might be caused by bandgap tailing due to a high con-centration of states in the gap near the band edges or to electric field effects on thebandgap due to these or other states. The larger bandgap of the nonoxidant filmscould be due to size quantization (the crystal size—as opposed to particle size—was not measured), band filling (Burstein–Moss shift), or, as suggested in the pa-per, a high concentration of oxygen in the PbS (although, as pointed out in Sec.5.2.3.3, the use of oxidant in the deposition solution did not affect the oxygen con-centration in the film).

Most other studies show the optical spectra of thin (often �100-nm) films,with emphasis on their solar control properties, and limited to a long-wavelengthlimit of 2.5 �m [46–48] (Ref. 49 shows reflection spectra to longer wavelengths).These films usually have an apparent absorption onset in the region of 600–800nm, the longer wavelengths characteristic of thicker films. The spectra at longerwavelengths are typically dominated by reflection rather than by absorption. Thiswould suggest bandgaps of between 2 and 1.5 eV. Figure 5.3 shows typical ex-amples of the transmittance and specular reflectance of PbS films deposited from

FIG. 5.3 Optical transmittance (a) and near-normal specular reflectance (b) of CD PbSfilms of different thicknesses. The thickness increases from A to F over an estimated rangeof ca. 50 nm to �200 nm. (Adapted from Ref. 46 with permission from IOP PublishingLtd.).


a Pb(Ac)2/triethanolamine/NaOH/thiourea bath; the absence of absorption in thenear-IR region, where PbS would normally absorb, is evident from a comparisonof the transmission and reflection data. In fact, observation of various CD PbS(from thiourea baths) spectra shows that there is absorption starting around 1.5–2�m, but it is usually weak. Partly this is due to the fact that much of the literaturedeals with thin films (often �100 nm), but the absorption does seem to be smallerthan expected. In the study by Pop et al. [49], moderately strong absorption stilloccurs in multilayer films (some hundreds of nanometers thick) at wavelengthslonger than 2 �m.

The small amount of information on films deposited from thiosulphate bathssuggest that the absorption of these films in the near-IR range may be higher thanthose deposited from thiourea baths. Gadave et al. measured a transmission under1% over the entire measured range from 2 �m to 400 nm (in fact the transmissionincreased slightly towards shorter wavelengths) [10]. No reflectivity or thicknessdata were given: however, the films must have been very thick (by CD standards)to give such a low transmission at 2 �m. Reference 17 shows a gradually de-creasing transmission from the maximum wavelength measured (2.5 �m) for afilm of unknown thickness, with a transmission of ca. 20% at 800 nm, while a 60-nm film deposited by a similar process on ultrathin Cu2S (which, according to thedata of Ref. 17, did not much affect the spectrum) showed a transmission of 11%at 800 nm (lower than most layers of comparable thickness deposited from athiourea bath) but with a poorly defined long-wavelength onset.

To sum up, while there is too little information available to draw any firmconclusions, it appears that films deposited from most thiourea baths are weaklyabsorbing in the near-IR region and that films deposited from thiosulphate solu-tions (which are mildly acidic) may possess different optical properties in generalthan those deposited from (alkaline) thiourea baths. In this respect, and if this dif-ference is real, it would be interesting to deposit PbS from thioacetamide baths,which can be both acidic and alkaline.

5.2.6 Photoconductivity of CD PbS (and PbSe):General Considerations

Specific photoconductive properties of PbS films have already been treated. Thissection deals with more general aspects of CD PbS (for the most part, also rele-vant for PbSe) films.

The bandgap of PbS at room temperature is 0.4 eV, corresponding to awavelength of 3.1 �m, which more or less gives the long-wavelength detectionlimit. The bandgap increases with increasing temperature, in contrast to the nor-mal semiconductor bandgap dependence on temperature. Therefore the long-wavelength detection limit of detectors made using these films shifts to shorterwavelengths with increasing temperature. A low temperature of operation there-


fore not only increases the sensitivity of the detectors (reduction of noise) but alsowidens the spectral response.

The air-annealing (typically between 300 and 400°C) of PbS and PbSe used,in most cases, to maximize the photoconductive response results in some surfaceoxidation, probably to PbSO3 and PbSO4. Such compounds have been detected inPbS precipitate, prepared from a typical CD process, after air-annealing at 300°Cor higher [50]. Oxidizing agents are often added to the CD bath to give photocon-ductive films in the as-deposited state [29,30]. In fact, commercial photoconduc-tive films used a proprietary oxidant as a matter of course. These commercialfilms, prepared from alkaline solutions of a lead salt and thiourea, were made upof several layers of CD PbS (total thickness ca. 1 �m) to maximize the photore-sponse, in particular the long-wavelength response.

The cause of photoconductivity in these films has been thoroughly studied.Here only a very condensed account of the theory is given. References 51 and 52give more detail.

Electrical conductivity, �, is given by the product of free-carrier concentra-tion, n, carrier mobility, �, and carrier charge, e:

� � ne� (5.5)

Therefore an increase in conductivity upon illumination (photoconductivity) canbe due to either an increase in carrier concentration and/or an increase in mobil-ity. In general, it is believed that an increase in carrier (hole) concentration is thedominant cause for room-temperature photoconductivity for the lead chalco-genides and that an increase in mobility becomes increasingly important at lowtemperatures. The dark conductivity of films deposited with or without added ox-idant were similar; the difference in photoconductivity between them was as-cribed to the formation of sensitizing centers (interband states) due to the oxidant.

Finally, it is worth mentioning a comment made in a paper describing junc-tions between Ge and CD PbS [34]. It was noted that evaporated epitaxial PbSfilms were poorly, if at all, photoconducting, while CD films, with mobility lowerby two orders of magnitude and much “poorer” structure, were much superior inthis respect. In a way, this should not be surprising since, for good photoconduc-tivity, low dark conductivity (and therefore either low mobility and/or low carrierconcentration) is necessary.

5.3 PbSe

PbSe has a considerably more recent history than does PbS; having been depositedby CD only about 60 years ago, instead of more than 130 years for PbS. The ear-liest work appears to have been carried out by the Germans during World War IIand, as for PbS, was shrouded in secrecy. While PbSe has been deposited by theSe analogues of thiosulphate and thiourea, as for PbS, in contrast to the history of


PbS deposition, the early work used mostly selenourea (or selenourea deriva-tives), while the more recent studies used the selenosulphate baths. As for PbS,these two baths will be treated separately. Optical and photoconducting propertieswill not be treated separately as for PbS, for which there is a larger body of results,but are included with the description of the various deposition baths.

5.3.1 Selenourea-Based Baths

The first apparent report in the open literature of CD PbSe for photoconductive de-tectors was in 1949 [53]. The PbSe was deposited from a solution of PbAc2 andselenourea onto a predeposited (from PbAc2 and thiourea) layer of PbS. The PbSlayer acted as a seed layer, presumably to obtain faster deposition (it was notedthat the PbSe deposition was much slower than that of PbS). The photoconduc-tivity of this film exhibited a broad maximum between 3 and 4 �m, giving a rea-sonable response out to beyond 4.5 �m (PbS drops off at 3 �m).

A detailed description of PbSe deposition using N,N-dimethylselenourea(DMS) was presented in 1964 [54]. DMS was used instead of selenourea becauseof its greater stability. Even so, Na2SO3 was added to inhibit decomposition to Se.The Pb (as nitrate) was complexed with citrate to keep it in solution in the alka-line conditions used (pH � 9.8 using ammonia). This pH was fairly critical: At 9.5(and at 25°C) the deposition was much slower, and at 10.1 rapid bulk precipita-tion occurred. Counter intuitively (and contrary to the case using selenosulphate),the reaction rate was faster for an aged DMS solution than for a fresh one (afteraging for 15 hr, little additional increase in rate occurred on further aging). Thiswas explained by inhibition of the DMS decomposition to selenide by the sulphite(Sec. 3.2.2.2). It was noted that PbSe formation occurred more rapidly in the ab-sence of sulphite and that adding fresh sulphite to an aged DMS solution reducedthe reaction rate. The best films were obtained if a fresh solution was used to de-posit PbSe on glass slides that had been precoated with a thin PbSe layer. In thiscase, film growth started almost immediately rather than after an induction period.

The resistivity of the films varied from �106 to �107 �. Since the thick-nesses of the films were not given, it is difficult to convert these values to a spe-cific resistivity. However, it does appear that the variations in resistivity were duemainly to film thickness and an upper value of 100 �-cm seems reasonable.

A detailed study of both the mechanism [55] and the kinetics [56] of PbSedeposition from selenourea baths has been carried out. The Pb was complexedwith citrate, and hydrazine was used to control alkalinity (remember that hy-drazine can function also as reducing agent and as complexant). As with the fore-going study, immediate deposition, with no induction time, occurred on glass onwhich PbSe had been precoated (or on glass that was sensitized with SnCl2 solu-tion, resulting in formation of tin hydroxide/oxide nuclei). On untreated glass, aninduction time for film deposition, which paralleled that for homogeneous pre-


cipitation in solution, was observed. The pH was somewhat lower than that of thepreceding study (typically ca. 8.8) and evidently less critical, although the depo-sition rate increased with increasing OH� concentration. The rate was also de-pendent on the selenourea concentration but independent of the Pb or citrate con-centration. Based on these observations, it was suggested that the rate-determiningstep was decomposition of selenourea by OH� to selenide ion:

(NH2)2CSe � OH� → HSe� � CH2N2 � H2O (5.6)

The selenide could then react with whatever Pb species were present, either in so-lution or as a solid phase, if present. The activation energy of this process, 60kJ/mole, was consistent with a chemical rate-determining step.

Films deposited from a selenourea bath (details not provided) were annealedat 350°C in air and changes in their resistivity and photoconductivity measured[57]. The resistivity of the films (between 1 and 4 �m thick) dropped from a fewhundred k� (assumed lateral conduction across the film) to about 1 k� after heat-ing for two minutes, with a peak in the resistivity after heating for about a minute.The photoconductivity maximum shifted from 1.5–1.8 �m for the as-depositedfilm to 3.6 �m for the annealed one. These phenomena were attributed to crystal-lization of a matrix of amorphous or nanocrystalline PbSe surrounding larger crys-tals, which crystallize or grow with annealing time. The crystal size of as-de-posited films in this and other studies by the same group was typically some tensof nanometers.

5.3.2 Selenosulphate-Based Baths

Sodium selenosulphate was used to precipitate PbSe from Pb2� solutions by mix-ing PbAc2 and Na2SeSO3 solutions [58]. The precipitation was rapid, and film formation did not occur (at least to any visible extent). Complexation was requiredto slow the reaction. Films of PbSe were first deposited using selenosulphate froma hydroxide-complexed (plumbite) bath [59]. By complexing the Pb [with citrate,Rochelle salt (tartrate), or oxalate] and adding alkali (NaOH, KOH, or NH4OH) toan optimum pH of 11.05, mirror films of PbSe were obtained [60]. The mecha-nism proposed, for both precipitate [58] and films using carboxylic acids [60], wasformation of PbSeSO3:

PbAc2 � Na2SeSO3 → PbSeSO3 � 2NaAc (5.7)

followed by decomposition of the PbSeSO3 to PbSe by water:

PbSeSO3 � H2O → PbSe � H2SO4 (5.8)

No evidence was presented for this mechanism, although it is a reasonable route,nor for the assumption of an identical mechanism for precipitate and film forma-tion (although that too is likely. In fact, in another study, based on a kinetic study


of PbSe formation from a similar bath employing glass powder (particle size ca.80 �m) as a “substrate,” it was concluded that the same mechanism was operativeboth for PbSe formation as a precipitate in solution and for that deposited on theglass surface [61]). On the other hand, Fofanov and Kitaev assumed a mechanismwhereby selenide ion, formed by hydrolysis of selenosulphate, reacted with Pb2�

via the ion-by-ion mechanism [59].The thickness of the films using the carboxylate complex was not dependent

on the Pb concentration (at least from 10 mM up). It was, however, very depen-dent on the pH value, with a sharp maximum, the position of which was depen-dent on the anion of the lead salt and, more so, on the hydroxide used (Na, K,NH4). The thickness varied from �50 nm to 300 nm. The film thickness was alsomoderately dependent on deposition temperature, increasing with increase in tem-perature (in contrast to the more usual decrease in terminal thickness with in-creasing temperature). It is not certain, however, whether these thicknesses (andothers measured in this study) were actually terminal thicknesses or only thethickness measured after a certain time.

The bandgap of these films was measured to be 0.28 eV—the same as theliterature value.

Films were also deposited from similar solutions, only using selenourea in-stead of selenosulphate and at somewhat lower pH (�9). The quality (of the mir-ror surface) and thickness of the films deposited from selenosulphate solutionswere similar to those deposited from the (at that time) conventional selenoureaones.

Kainthla et al. carried out an investigation on the parameters that affecteddeposition rate [62]. The rate increased, as expected, with increase in temperatureand selenosulphate concentration. However, it decreased with increase in pH. Thiswas explained on the basis of the expected dominant hydroxy-citrate complex ofPb, [Pb(OH)C6H5O7]2�. The equilibrium involving this complex is

[Pb(OH)C6H5O7]2�D Pb2� � OH� � C6H5O73� (5.9)

A greater hydroxide concentration will shift the equilibrium to the left, decreasingthe free-Pb2� concentration. This explanation means that the rate is dependent onthe lead concentration, in contrast to the previous studies discussed earlier, whichfound the rate to be essentially independent of this parameter.

It was also noted that the optimal pH was temperature dependent: At lowdeposition temperature it was �9, while at high temperatures it was �10. This fol-lows from the inverse dependence of rate on pH. The optimal pH is a balance be-tween slow-enough formation of PbSe (to prevent precipitation in solution) butnot too slow to prevent formation of a film in a reasonable time. At higher tem-peratures the rate is faster, and therefore the optimal pH should be lower.

Films (between 0.4 and 2.5 �m thick) deposited from selenosulphate so-lution were characterized by optical spectroscopy and (photo)conductivity [63].


There appeared to be a very large variation in optical properties from film tofilm; in one case, little absorption (corrected for specular reflection) was mea-sured from 2 �m toward increasing wavelength, while another sample showed asharp absorption onset close to 5 �m. The films had a resistivity of between3.103 and 3.104 �-cm and a photoconductivity maximum varying between 1.2and 1.9 �m.

An investigation of PbSe deposited from selenosulphate baths was carriedout using various complexants and conditions and with emphasis on the mecha-nisms and, in particular, on the crystal size and morphology of the deposits [64].Citrate, nitrilotriacetate (NTA), and hydroxide (in order of increasing complexingstrength with Pb) were used under conditions where a solid hydrated oxide phasewas either present in solution (cluster) or absent (ion-by-ion). The presence of thisphase, and a semiquantitative estimation of its relative concentration, was mea-sured by UV absorption and light-scattering measurements. As for CdSe, the mor-phologies of the two types of films were very different, with the ion-by-ion filmshaving a larger crystal size. However, the difference depended on the complexantused as well as on other deposition conditions. A wide variety of crystal sizes wasobtained. Figure 5.4 shows some TEM micrographs of various PbSe films. Theconditions of deposition for each film (the important factors are given in the fig-ure legend) are not important here; the purpose is to demonstrate the wide rangeof crystal sizes and morphologies in PbSe films obtained using different com-plexants, solution compositions, and deposition temperatures.

Under certain conditions, relatively large and small crystals coexisted in thesame films. For example, in “cluster” deposition from citrate complex, both small(ca. 4 nm) and medium-sized (6–12 nm) crystals were formed, although only thesmall crystals formed at the beginning of the deposition, and the larger ones ap-peared in thicker films, again, an indication of both mechanisms operating (ion-by-ion growth normally is slower than cluster growth, therefore it takes longer forthe larger crystals to appear). Additionally, the small and larger crystals were notdeposited together but in separate regions; i.e., regions only of larger crystals wereformed surrounded by (most of the deposit) only small crystals. It was suggestedthat the larger crystals formed in regions when no previous small-crystal deposi-tion had occurred (i.e., in pinholes in the originally thin film).

For low selenosulphate concentrations, only the small crystals were formed,even in thicker films, and this was rationalized by the lower steady-state selenideconcentration, which would favor cluster growth over ion-by-ion formation (theproduct of free lead and selenide ions needs to be larger than the solubility productof PbSe for ion-by-ion deposition to occur). An important difference between thecitrate depositions and the NTA or hydroxide ones is that, even in the “ion-by-ion” citrate deposition, some low concentration of colloidal hydrated oxide waspresent, due to the relatively low complexing strength of citrate. The pH of the hy-droxide baths (�13) was much higher than that of the citrate or NTA baths (10.8).


In common with CdSe deposition from selenosulphate baths, cluster depo-sition of PbSe normally resulted in specular films, while ion-by-ion films were ini-tially highly scattering as thin films but eventually (usually) became specularly re-flecting with increase in thickness. As for CdSe, the development of specularitywith thickness of ion-by-ion films could be explained by filling in of voids be-tween the large crystals.

The wide range of very small crystal sizes in these films gives rise to strongblue shifts in their optical spectra due to size quantization [65]. This aspect ofthese films is dealt with in detail in Chapter 10.

Other complexants have been used for PbSe deposition. Triethanolaminewas used in one study [66]. While deposition occurred over a wide range of tem-peratures, optimum results (in terms of rate of deposition and film thickness) wereobtained at a deposition temperature of 75°C. In another study, lead nitrate wasdissolved in an excess of hydroxide and excess selenosulphate was also used as anadditional complexant [67]. The pH was 10 (adjusted with acetic acid), and depo-

FIG. 5.4 TEM micrographs of PbSe films deposited under different conditions. A: fromcitrate/0°C/hydroxide mechanism. B: nitrilotriacetate/60°C/ion-by-ion mechanism. C: hy-droxide/60°C/ion-by-ion mechanism. D: hydroxide/80°C/ion-by-ion mechanism. All filmsare as deposited.


sition was carried out both at room temperature and at 70°C. The substrate (glassor polyester) was pretreated with SnCl2 solution—the standard tin sensitization;presumably this improved the adherence and/or homogeneity of the films. The re-sistivity of these films was ca. 104 �-cm as deposited. Annealing in air at 100°Cfor 24 hr decreased this resistance to 3 � 10�3 �-cm, a drastic change for such amild treatment and one that suggests that the electrical properties of these as-de-posited films may change greatly with time, even under ambient conditions.

5.3.3 Variations in PbSe Deposition

X-ray amorphous PbSe has been deposited by addition of Na2S2O3 (at least halfthe selenosulphate concentration) to a room-temperature selenosulphate/citratebath [68]. The thiosulphate increased the induction time for PbSe formation in thesolution, although eventually film growth was even faster than in the absence ofthiosuphate. It was noted that the initial, very thin deposit was yellow. A clearXRD pattern was obtained (for thicker films) after annealing at 350°C for oneminute, giving a crystal size of 13 nm (which grew with continuing annealing). Itis interesting to note that, in contrast to films deposited in the absence of thiosul-phate, those made using thiosulphate did not peel off, even after remaining a longtime in the depositing solution. Chemically deposited films will often peel off ifleft too long in solution, probably due to stress in the thickening film, and this isexpected to be less or absent in an amorphous film. The amorphous structure wasattributed mainly to the increased deposition rate caused by the thiosulphate.Since the growth rate was apparently in the range of hours, this by itself would notbe expected to result in an amorphous film, although it might contribute to the ef-fect. A thorough investigation of this deposition would be desirable, includingTEM imaging (to see whether the film is, indeed, amorphous, since XRD by itselfis not a good enough verification of this) and, particularly important, composi-tional analysis (to see if there is an appreciable amount of S in the film, thiosul-phate itself forming PbS).

From the optical spectrum, an approximate bandgap of 1.5 eV could be esti-mated, although the interpretation of the spectrum is open to ambiguity. A photo-conductivity maximum at 1.1 �m (1.1 eV) was measured. These anomalously highenergy values were attributed to size quantization (Chap. 10). Transmission in themid-IR was also measured for these films. Although there were some absorptionbands, probably due to adsorbed species from the deposition bath, the films wereessentially transparent at wavelengths longer than ca. 5 �m, the transmission de-creasing gradually at shorter wavelengths. While annealing at 350°C did changethe IR spectral shape, particularly the intensity of some of the absorption bands,overall the transparency beyond 8 �m was maintained (in some regions, even in-creased). This implies that the films, both as deposited and annealed, were very in-trinsic, since free carriers would result in absorption in this spectral region.


The (presumed lateral) resistance of the films was ca. 108 �, which de-creased 4–5 orders of magnitude after annealing at 350°C.

In analogy with PbS, a film of PbSe has been formed by passing H2Se overPbAc2 solution acidified with acetic acid, giving specularly reflecting films thatcould be picked up from the surface of the solution [25]. Electron diffraction ofthese films showed a tendency to grow with (001) planes parallel to the solutionsurface. The average crystal size of these films was 25 nm.

A variation of the CD process for PbSe involved deposition of a basic leadcarbonate followed by selenization of this film with selenosulphate [64]. Whitefilms of what was identified by XRD as 6PbCO3�3Pb(OH)2�PbO (denoted here asPbMOHMC) were slowly formed over a few days from selenosulphate-free so-lutions that contained a colloidal phase and that were open to air (they did not formin closed, degassed solutions). CO2 was necessary for film formation—other thansparse deposits, no film formation occurred of hydrated lead oxide under any con-ditions attempted in this study. Treatment of these films with selenosulphate so-lution resulted in complete conversion to PbSe at room temperature after 6 min.The selenization process of this film was followed by XRD, and it was seen to pro-ceed by a breakdown of the large PbMOHMC crystals to an essentially amor-phous phase of PbSe with crystallization of this phase to give finally large (ca. 200nm) PbSe crystals covered with smaller (15–20 nm) ones as well as some amor-phous material.

Hydrazine added to a selenourea//Pb2� bath was found to strongly affect theelectrical and photoconductive properties of the PbSe films [69]. As the hydrazineconcentration increased, the dark resistivity increased in a very non-monotonicway, from ca. 0.1 �-cm (low hydrazine concentration) to 104 �-cm. Photocon-ductivity was observed, as might be expected, only for the high-resistivity films.The effect of the hydrazine was attributed to an increase in the deposition rate byan increase in pH; however, the pH values of the solutions were not reported, andhydrazine would be expected to affect the reaction as a reducing agent apart frompH considerations. As an interesting aside, it was noted that only p-type filmscould be deposited by this technique (and this is the case for most CD PbS), whileevaporated PbSe films were normally n-type (due to Se vacancies). The absorp-tion edge of the films was ca. 4.2 �m, close to to the literature bandgap of 0.28 eV.

5.3.4 Comparison of Films Deposited fromSelenourea and Selenosulphate Baths

A series of studies comparing films deposited from selenourea and selenosulphatebaths were carried out, with emphasis on photoconductivity behavior and the ef-fects of annealing. A broad photoconductivity maximum occurred at ca. 1.6 �mfor the selenosulphate film, decreasing strongly beyond 2 �m. A similar spectrum,


but shifted 0.2–0.4 �m to longer wavelengths, was observed for the selenoureafilm. Annealing at 350°C in air broadened both spectra to ca. 5 �m (approximatelythe bulk bandgap of PbSe), although the the photoconductivity response of the se-lenosulphate films was initially degraded (mainly due to increased noise), and along (ca. 50 hr) annealing time was required for good response [70].

Photoluminescence spectra of the films were measured (77 K) and com-pared with an epitaxial PbSe layer in the same study. Blue shifts in the spectra(greater for the selenourea films) were attributed to quantum size effects (seeChap. 10). The crystal size was reported to be 40–60 nm (the selenourea ones be-ing somewhat smaller than the selenosulphate ones), growing to 100–150 nm af-ter annealing.

The changes in resistivity with annealing of films deposited from selenoureaand selenosulphate baths, as well as evaporated films, were compared [71,72]. Al-though there were small differences between the various films, no major differ-ence was found. Additionally, the resistivity of as-deposited films, deposited fromboth selenourea and selenosulphate baths, does not change with time over a periodof months in air. However, after annealing in air at 350°C when the resistivity in-creases, there is a gradual decrease in room-temperature resistivity (and also inphotoconductivity response) with time [73]. These variations were related to for-mation of PbSeO3 and adsorbed oxygen on the surface of the annealed crystals.

REFERENCES1. C Puscher. Dingl. J. 190:421, 1869.2. J Emerson-Reynolds. J. Chem. Soc. 45:162, 1884.3. DE Bode. In: Physics of Thin Films. Vol. 3, Academic Press, New York, NY, 1996,

p 275.4. F Kicinski. Chem. Ind. 54, 1948.5. TS Moss. Proc. IRE 43:1869, 1955.6. WN Arnquist. Proc. IRE 47:1420, 1959.7. DJ Lovell. Am. J. Phys. 37:467, 1969.8. TH Johnson. Proc. SPIE 60, 1983.9. CD Lokhande. Mater. Chem. Phys. 28:1, 1991.

10. KM Gadave, SA Jodgudri, CD Lokhande. Thin Solid Films 245:7, 1994.11. E Beutel, A Kutzelnigg. Z. Elektrochem. 36:523, 1930.12. PL Allen, A Hickling. Chem. Ind. 51:1558, 1954.13. EB Andersen. Z. Phys. Chem. B. 32:237, 1936.14. WH Perkins, AT King. J. Chem. Soc. 103:301, 1913.15. E Beutel. Z. Angew. Chem. 26:700, 1913.16. E Beutel, A Kutzelnigg. Monats. 58:295, 1931.17. I Grozdanov. Semicond. Sci. Tech. 9:1234, 1994.18. M Najdoski, B Mincevasukarova, A Drake, I Grozdanov, CJ Chunnilall. J. Mol.

Struct. 349:85, 1995.19. HL Smith. J. Sci. Instrum. 4:115, 1927.


20. O Hauser, E Biesalski. Chem. Ztg. 1078, 1910.21. G Bruckman. Kolloid Z. 65:1, 1933.22. MK Norr. J. Phys. Chem. 65:1278, 1961.23. H Pick. Zeit. Phys. 126:12, 1949.24. HN Acharya, NK Misra. J. Phys. D: Appl. Phys. 4:1968, 1971.25. H Wilman. Proc. Phys. Soc. 60:117, 1948.26. HN Acharya, HN Bose. Phys. Status Solidi (a) 6:K43, 1971.27. VF Markov, LN Maskaeva, GA Kitaev. Inorg. Mater. 36:657, 2000.28. EM Larramendi, O Calzadilla, A Gonzalez-Arias, E Hernandez, J Ruiz-Garcia. Thin

Solid Films 389:301, 2001.29. GH Blount, PJ Schreiber, DK Smith, RT Yamada. J. Appl. Phys. 44:978, 1973.30. GP Kothiyal, B Ghosh, RY Deshpande. J. Phys. D: Appl. Phys. 13:869, 1980.31. IC Torriani, M Tomyiama, S Bilac, GB Rego, JI Cisneros, ZP Argüello. Thin Solid

Films 77:347, 1981.32. PK Nair, VM Garcia, AB Hernandez, MTS Nair. J. Phys. D: Appl Phys. 24:1466,

1991.33. J Bloem. Appl. Sci. Res. Section B 6:92, 1956.34. JL Davis, MK Norr. J. Appl. Phys. 37:1670, 1966.35. G Guizzetti, F Filippini, E Reguzzoni, G Samoggia. Phys. Status Solidi (a) 6:605,

1971.36. S Watanabe, Y Mita. J. Electrochem. Soc. 116:989, 1969.37. M Isshiki, T Endo, K Masumoto, Y Usui. J. Electrochem. Soc. 137:2697, 1990.38. BL Sharma, SN Mukerjee. Phys. Status Solidi (a) 2:K21, 1970.39. H Sigmund, K Berchtold. Phys. Status Solidi 20:255, 1967.40. FC Meldrum, J Flath, W Knoll. J. Mater. Chem. 9:711, 1999.41. FC Meldrum, J Flath, W Knoll. Thin Solid Films 348:188, 1999.42. I Pintilie, E Pentia, L Pintilie, D Petre, C Constantin, T Botila. J. Appl. Phys. 78:1713,

1995.43. S Gallardo, M Gutierrez, A Henglein, E Janata. Ber. Bungenges. Phys. Chem.

93:1080, 1989.44. AF Gibson. Proc. Phys. Soc. B 63:756, 1950.45. JGN Braithwaite. J. Sci. Instr. 32:10, 1955.46. PK Nair, MTS Nair, A Fernandez, M Ocampo. J. Phys. D: Appl. Phys. 22:829, 1989.47. PK Nair, MTS Nair. J. Phys. D: Appl. Phys. 23:150, 1990.48. PK Nair, M Ocampo, A Fernandez, MTS Nair. Sol. Energy Mater. 20:235, 1990.49. I Pop, C Nascu, V Ionescu, E Indrea, I Bratu. Thin Solid Films 240: 1997.50. GA Kitaev, LG Protasova, VG Kosenko, VF Markov, LV Dyadyuk, GV Stepanova.

Inorg. Mater. 29:1182, 1993.51. S Espevik, C-H Wu, RH Bube. J. Appl. Phys. 42:3513, 1971.52. E-H Lee, RH Bube. J. Appl. Phys. 43:4259, 1972.53. CJ Milner, BN Watts. Nature 163:322, 1949.54. RA Zingaro, DO Skovlin. J. Electrochem. Soc. 111:42, 1964.55. AB Lundin, GA Kitaev. Inorg. Mater 1:1900, 1965.56. AB Lundin, GA Kitaev. Inorg. Mater. 1:1905, 1965.57. LP Biró, A Darabont, P Fitori. Europhys. Lett. 4:691, 1987.58. NI Glistenko, AA Eremina. Zh. Neorg. Khim. 5:1003, 1960.


59. GM Fofanov, GA Kitaev. Russ. J. Inorg. Chem. 14:322, 1969.60. NI Glistenko, IN Shramchenko, AN Kosilova. Zhurn. Prikl. Khim. 45:1356, 1972.61. GA Kitaev, AZ Khvorenkova. Russ. J. Appl. Chem. 72:1520, 1999.62. RC Kainthla, DK Pandya, KL Chopra. J. Electrochem. Soc. 127:277, 1980.63. RM Candea, N Dadarlat, R Turcu, E Indrea. Phys. Stat. Sol. (a) 90:K91, 1985.64. S Gorer, A Albu-Yaron, G Hodes. Chem. Mater. 7:1243, 1995.65. S Gorer, A Albu-Yaron, G Hodes. J. Phys. Chem. 99:16442, 1995.66. P Pramanik, S Biswas, PK Basu, A Mondal. J. Mater. Sci. Lett. 9:1120, 1990.67. I Grozdanov, M Najdoski, SK Dey. Mater. Lett. 38:28, 1999.68. LP Biro, RM Candea, G Borodi, A Darabont, P Fitori, I Bratu, D Dadarlat. Thin Solid

Films 165:303, 1988.69. DH Roberts, JE Baines. J. Phys. Chem. Solids 6:184, 1958.70. D Dadarlat, RM Candea, R Turcu, LP Biro, Zasavitskii, II, MV Valeiko, AP Shotov.

Phys. Status Solidi (a). 108:637, 1988.71. RM Candea, R Turcu, P Margineanu, N Dadarlat. Phys. Status Solidi (a) 96:337,

1986.72. RM Candea, R Turcu, G Borodi, I Bratu. Phys. Status Solidi (a) 100:149, 1987.73. RM Candea, LP Biro, N Dadarlat, G Borodi, A Darabont, P Fitori, R Turcu. Phys.

Status Solidi (a). 108:233, 1988.


6Other Sulphides and Selenides

6.1 INTRODUCTION

This chapter covers sulphides and selenides not included in Chapters 4 and 5, i.e.,all metals except for Zn, Cd, Hg, and Pb. Some of these materials, e.g., the sul-phides of Bi, Cu, and Ag and Cu-Se, have been the subject of many investigations.There are others, however, on which as little as one paper has been published al-together.

In order to minimize, as much as possible, making this chapter into a list ofrecipes and properties, its layout will be somewhat different than that used up tonow. The chapter is divided into sections alphabetically (by English name ratherthan by chemical symbol), designated by the name of the metal. Basic data for thosecompounds for which there are at least several papers will be listed in tabular formtogether with references. The data in the first column of the table (the depositionbath composition) will be given as sulphide or selenide source/complex/tempera-ture and pH. The specific metal ion is not given in most cases, because this will beclear from the location of the table and the metal salt used will normally be foundin the specific description of the study in question. It should be kept in mind thatthe chalcogen source itself is a complexant, often fairly weak but sometimes strong.Resistivity is given as specific resistivity; if not defined, then the resistance re-ported is given in parentheses. Specific details outside of these basic data will betreated separately in the order in which the particular study appears in the table and,


if connected with a specific datum in the relevant table, will be signified by an as-terix after this datum. Each discussion of this type will include the reference num-ber to identify it. Abbreviations used in the “Solution” column are: DMSeU—N,N-dimethyl selenourea; EDTA—ethylenediaminetetra-acetate; S2O3—thiosulphate;SeSO3—selenosulphate; TA—thioacetamide; TEA—triethanolamine; TU—thiourea; RT—room temperature. The abbreviations in the “Bandgap” column are:dir.—direct; indir.—indirect.

Before discussing the specific studies, a general overview will be given foreach metal. This overview relates only to information of direct relevance to thesubject and will be relatively lengthy in some cases and very brief or even nonex-istent in others, according to need.

6.2 ANTIMONY

6.2.1 Sb2S3 (See Table 6.1)

All reported cases of CD antimony sulphide involve the trisulphide, Sb2S3. Sb2S3

is soluble in hydroxide to give antimonates and in excess sulphide to give thio-complexes. The latter is not a problem in CD since free sulphide, if it exists, doesso in a very low concentration and is rapidly taken up to precipitate Sb2S3. How-ever, the solubility in alkaline solutions limits the pH of the deposition solutions.Sb2S3 exists in two forms: so-called amorphous Sb2S3, which varies in color from

TABLE 6.1 Antimony Sulphide

Bandgap Resistivity Conduct.Solution (eV) (�-cm) type Ref.

TA/NH3/TEA/RT 1.85 ind. 5 � 104 n- 3Annealed, 300�C 1.74 ind. 250TA/NH3/TEA/RT 1.86 ind. 4 � 108 n- 4Annealed, 300�C 1.74 ind. 5 � 106 (5 � 108) 5TA/tartrate/RT/pH 9.5 1.62 ind. 107 6TA/SbCl3 in CH3COOH 1.75 dir. 7 � 105 (130�C) n- 7

(nonaqueous)/RT*/pH 1.5Annealed 200�C 1.62 dir. 2 � 105 (130�C) n-TU/RT/pH 1–1.2 8Annealed � 200�C 108 n-S2O3/SbCl3/EDTA/RT/pH 2–3 1.82 104–105 9S2O3SbCl3 in CH3 COOH 2.48* dir. �107 n- 10, 11Annealed 170–200�C 1.76 dir. �105

S2O3/SbCl3 in acetone/10�C/pH 5 2.21 dir. �107 12Annealed 250�C 1.79 dir.


yellow to red, and black crystalline Sb2S3, with a bandgap between 1.7 and 1.8 eV.Thioacetamide is the most commonly used S source for this compound.

Antimony sulphide deposition onto metallic substrates, together with PbSand Cu-S, was first reported in the original paper of Puscher [1] using thiosulphateand antimony tartrate. No characterization of this film was carried out nor prop-erties given. The same method was also described “recently,” in 1931, using anumber of different metals as substrates [2]. It was noted that SbCl3, when mixedwith thiosulphate, reacted too rapidly, hence the use of the tartrate (tartaric acid isa complexant). Again, no characterization of the films was made.


In Refs. 3–5, potassium antimonyl tartrate was used as a source of Sb. While thedeposition conditions appeared to be almost identical in all cases (other than thesilicotungstic acid added in some cases in the latter two studies), there were someunexplained differences in structural and electrical properties (the latter shown inthe table). The lower resistivities obtained in the first study were mirrored byhigher carrier concentration (2 � 1015 cm�3 vs. 1012 cm�3) and slightly highermobility (14 vs. 10 cm2V�1 sec�1).

The as-deposited film gave either no peaks or very broad ones in the XRDpattern, from which it is difficult to extract a crystal size, since there are so manyclose peaks in the spectrum. After annealing, well-defined peaks were observed.The films in the first study (deposited on glass) had a grain size (measured bySEM, not necessarily crystal size) of 0.12 �m as deposited, growing to nearly 4 �m after annealing. From Ref. 5, the grain size of the annealed films on SnO2-coated glass or steel (both with well-defined crystal surfaces) was close to 1 �m,but for films deposited on glass it was 0.08 �m. The films in the first study werereported to be close to stoichiometric, while those in the other studies were moreor less S rich.

Addition of silicotungstic acid to the bath [4,5] reportedly improved the sto-ichiometry and increased the grain size somewhat (as well as introducing a sepa-rate WO3 phase) of the annealed films. There was no major effect of the STA onthe electrical resistance. The resistivity of the annealed (not the as-deposited)films was also very substrate dependent [5]: ca. 5 � 108 �-cm for films on glassor steel and two orders of magnitude less on SnO2-coated glass. Therefore there isno direct relationship between the resistivity and the grain size. Photoconductiv-ity in these layers was studied [5]. The ratio between dark and light resistivitieswas as high as 104 (for annealed films deposited with silicotungstic acid), but thisratio decreased, and the decay time increased, with exposure to air. The ratio foras-deposited films was ca. 10. Photoconductivity decay times varied from �1 secto several seconds.

In a similar study [6], SbCl3 was complexed with tartaric acid. Also,NaOH was used instead of ammonia, and the solutions were more concentrated


than the previous one. The films were close to stoichiometric. A lower value ofindirect bandgap was measured for this study (1.62 vs. 1.85 eV). It is notewor-thy that while the films exhibited no XRD pattern, the precipitate in solutiongave sharp XRD peaks. Whether this is due to a different mechanism for the tworoutes or an effect of the glass substrate, two suggestions made by the authors,is an open question.

A nonaqueous (acetic acid) bath is described in Ref. 7. While the depositionwas carried out mostly at room temperature, the bath was initially heated to 80°Cfor 10 min to speed up the deposition. Unlike other films of Sb2S3, the as-deposited films showed an XRD pattern of well-defined peaks (although these be-came more numerous and with mostly different reflections on annealing). It is in-teresting to note, in possible connection with this observation, that the as-de-posited films were described as pink, which is not typical of any form of Sb-S andsuggests a mixture of a white material with a reddish one.

6.2.1.2 Thiourea

Reference 8 is the only study in which thiourea has been used to deposit Sb2S3. Amethanolic solution of SbOCl at a very low pH (ca. 1) was used. This seems to bea unique case of deposition using thiourea at low pH and suggests that the reactionmay proceed through some complex-decomposition reaction, since free sulphideis not expected to form under such conditions. The as-deposited films were highlyscattering, white, and nonadherent, but converted to adherent, still scattering filmswith the typical orange color of as-deposited Sb2S3 after heating at 120°C. X-raydiffraction of this film showed no pattern; annealing at over 200°C converted thefilm to gray Sb2S3 with a well-defined XRD spectrum. While optical spectra weregiven, it is difficult to interpret them, due to the large degree of scattering.


Reference 9 involved the first thiosulphate deposition after the original reports of Refs. 1 and 2 and gives some properties of the films but no structural charac-terization.

In Refs. 10 and 11, aqueous Na2S2O3 was added to SbCl3 in glacial aceticacid (SbCl3 hydrolyzes in water unless complexed or the solution is moderatelyacidic or strongly alkaline). A pH of ca. 3 was optimum; below 2.5, adhesionwas poor; above 4, basic antimony salts precipitated. The solution was kept be-low room temperature to prevent rapid bulk precipitation. No XRD pattern wasfound for the as-deposited film, which was presumed to be amorphous. Anneal-ing at 170°C crystallized the film, at least partly. The bandgap of the as-de-posited film was reported to be 2.48 eV and that of the annealed film 1.76 eV.Photoconductivity was exhibited by the annealed film but not by the as-de-posited one.


A similar thiosulphate bath, only using acetone instead of acetic acid to dis-solve the SbCl3 was described in Ref. 12. It was reported that the films made us-ing acetic acid tended to flake off, and use of acetone prevented this. Photocon-ductivity studies showed that the photosensitivity was poor for as-deposited films(ca. a factor of 2 increase in conductivity) but increased to between 102 and 103

after optimum annealing at 300°C in N2 (annealing at 250°C or less had little ef-fect on the photosensitivity).

6.2.2 Sb2Se3

Only one group has reported CD of Sb2Se3. The solution used was potassium an-timonyl tartrate, complexed with triethanolamine and ammonia. Selenosulphatewas used as the Se source. No XRD pattern was found, as for the sulphide de-posited under equivalent conditions. The bandgap was 1.88 eV, and resistivity�107 �-cm [13,14]. Continued study of this deposition showed the effect of var-ious parameters on deposition rate and film thickness (the latter typically reaching1 �m) [15]. This study also described some photoelectrochemical behavior ofthese films (Chap. 9).

6.3 ARSENIC

There are at least three sulphides of As: As4S4, As4S6, and As4S10. The dimeric ar-senous (III) sulphide, As4S6 (often given as As2S3) and arsenic (V) pentasulphide(As4S10) can be precipitated by H2S as yellow solids from acidic solutions of therespective As salts. These sulphides are soluble in alkaline solution and in(poly)sulphide solutions and must therefore be deposited from acidic or at mostneutral solution. The pentasulphide is not a very stable compound; it is hydrolyzedby boiling water to arsenious acid and also is unstable as a solid in air above ca. 100°C.

Only two papers on CD of an arsenic chalcogenide (arsenic sulphide) werefound. Films were obtained using thiosulphate in an EDTA solution of As2O3 atroom temperature and a pH of 2–3 [9]. No compositional information was given.A bandgap of 2.0 eV and a resistivity of 104–105 �-cm were measured.

As2S3 was deposited at room temperature (27°C) from an acidic (pH � 2)thioacetamide bath containing As2O3 dissolved in concentrated HCl (and in somecases complexed with EDTA) [16]. The terminal thickness (which reached a max-imum and then decreased with time) was studied as a function of various deposi-tion parameters. Well-defined XRD peaks were obtained showing the monoclinicstructure (notable since this compound has a tendency to be amorphous or nearlyso as deposited). A direct bandgap of 2.42 eV (similar to the standard value forAs2S3) was estimated from the optical spectrum. The resistivity was ca. 105 �-cm.


6.4 BISMUTH

6.4.1 Bi2S3 (See Table 6.2)

Bismuth sulphide, Bi2S3, has been rather extensively investigated. Bi3� is readilyhydrolyzed in aqueous solution and is either used in acid solution or strongly com-plexed. Note the very low-solubility product of this compound—10�98 (Table1.1). The very low value is due largely to the large number of ions involved (five).However, even apart from this, the solubility (given as the concentration of freeBi3� and S2� ions) is low, and the solid is very readily precipitated.

A glance at the various values of bandgap shown in the table is enough tosee how wide this range is. These differences are often attributed either to anamorphous structure or to size quantization. However, as discussed in more detailin Chapter 10, many of these values are not dependable, and the values given inthis chapter are, in many cases, more a measure of the shape of the transmission(absorption) curve than of any specific bandgap value.

Deposition of Bi2S3 was first reported in 1931 from a thiosulphate bath ontometal substrates, although no details of the deposit properties were provided [2].


In Ref. 18, the initial deposition temperature was 60°C, which was then loweredto 27°C. The properties of the films were found to vary considerably with filmthickness. The crystal sizes and bandgaps are discussed in Chapter 10. Films 220

TABLE 6.2 Bismuth Sulphide


S2O3/EDTA/60–27�C*/pH 2 2.22–1.62* dir. 5.103–106 9, 17, 18S2O3/HCHO/CH3COOH/ 1.9 dir. 106 n- 19

7�C/pH 1.4S2O3/HCHO/CH3COOH/ 1.27 dir. 105 20

25�C/pH 1.5TA/TEA/NH3 1.7 indir. 103 p- 21TA/TEA/25 or 50�C/pH 8.5 1.9 ca. 106* 22, 23

1. 6 (illum.)TA/(�EDTA)/pH1–2 1.84 dir. 105 24TU/TEA/NH3/ ca. 1.5* 107 �-cm n- 25, 26

100�C-�RT*/pH 8–10TU/TEA/NH3/95�C/pH 9.5 1.6 105–106 n- 27TU/NH3/up to 90�C 1.76 28


nm thick had a resistivity of 5 � 103 �-cm, compared to 53-nm-thick films witha value of ca. 3 � 104 �-cm.

Reference 19 described a nonaqueous deposition using a mixture of aceticacid (in which the Bi(NO3)3 was dissolved) and formaldehyde (in which thesodium thiosulphate, which was insoluble in acetic acid, was dissolved). Deposi-tion was carried out at 7°C because the resulting films were more homogeneousthan at higher temperatures. No clear XRD pattern was found for the films, al-though one was for the precipitated powder.

A more recent investigation of the previous deposition [24] showed an XRDpattern of moderately sharp peaks similar (in peak widths) to the precipitated pow-der but weaker (presumably due to less material) and with only some of the re-flections (suggesting texturing of the film).


The majority of the various studies on Bi2S3 used thioacetamide as the source ofsulphur. In the first of these studies [20], very broad peaks were observed in theXRD spectrum and the film was classified as amorphous. (A radial distributionfunction analysis of these films allowed a structure to be proposed [29]). Muchsharper XRD peaks were obtained after a mild annealing at 150°C (for 6 hr). Thebroad peaks in the as-deposited film seem, for the most part, to have different values of 2� from those of the annealed film.

In Ref. 21, it was noted that use of ammonia (or NaOH) resulted in filmscontaining particulate deposits and incorporation of Bi(OH)3 into the films.Therefore the deposition solution was ammonia free, and, although alkaline, thepH was relatively low. It was also noted that use of thiourea instead of thioac-etamide gave very thin films. No difference in either optical or electrical proper-ties was found for films deposited at 25 or 50°C. The dark resistivity dropped fromca. 106 �-cm to as low as 0.3 �-cm (this value was apparently very variable fromsample to sample) after air-annealing at 200°C, although it increased steeply againat higher annealing temperatures. The photoconductivity of these films was stud-ied. The sensitivity was ca. 20 for the as-deposited films, increasing to ca. 100 ormore after (optimum) annealing at 150°C. This temperature corresponded to thetemperature at which the dark resistivity began to decrease strongly; above thistemperature, the dark resistance dropped and therefore so did the photoconduc-tivity response. The effect of various annealing treatments on the films was stud-ied in more detail in a separate work [30]. Here, XRD showed no peaks whatso-ever for the as-deposited films but strong and sharp peaks after annealing at200°C. Specific studies of the effect of annealing on the electrical and photocon-ducting properties of these films annealed in argon or hydrogen [31] and in oxy-gen [32] have been described but will not be discussed further here.

The rate of deposition of these films increased on exposure to light. Thisphenomenon is often observed in CD films and is believed to be due to photogen-


erated charge formation in the growing particles, leading to an electrochemical de-position process in parallel with the CD one. This phenomenon has been exploitedto generate photographic images by shining light through a mask (in this case, aphotographic negative) onto the depositing film [22]. The films in this case weredeposited on a predeposit of ZnS to improve adhesion. The contrast (mainly intransmission but also seen in reflection) between the thinner and thicker regionsof the film reproduced the image on the negative. Changes in bandgap due to sizequantization may also have contributed to the contrast (see Chap. 10).

It was mentioned that better adhesion was obtained on glass if ZnS was pre-deposited first [22,33]. Improved adhesion was also obtained from films depositedfrom this bath if the glass substrates were first treated with an organosilane [22].The silane binds to the glass surface, and the growing film anchors to terminalamino or thiol groups, both of which (the thiol in particular) bind strongly to metalchalcogenides.

A fairly strongly acidic thioacetamide bath (pH between 1 and 2) was de-scribed in Ref. 23, both with and without EDTA (sodium salt) as a complexantWhile some structural differences as well as variations in film thickness werenoted between the films deposited from baths with and without EDTA, they werenot highly significant. As with most other acidic baths, the films were clearly crys-talline, showing defined XRD peaks that allowed a crystal size of the order of 10nm to be estimated.

6.4.1.3 Thiourea

Triethanolamine was used in Ref. 25 as complexant together with ammonia, thelatter to slow the reaction, and it also apparently improved adherence. The solu-tion was heated to boiling for 40 min and then left at room temperature for 4 hr.As for other depositions using this strategy, no rationale was given for this regime,although some hints as to its reason might be gleaned from the follow-up paper([26]; see next paragraph): It can be presumed that it provided better films thanthose obtained by simply depositing at one temperature. No XRD pattern was ob-tained as for other alkaline bath depositions.

In a follow-up investigation [26], the initial pH was 10.17, and this droppedsomewhat during the deposition. The deposition was dependent on the age of theBi(NO3)2/triethanolamine solution. A solution aged for ca. 5 hr before making upthe deposition solution gave films about twice as thick (0.1–0.2 �m) as a freshlymade one. However, if aged for 24 hr, the color of the heated solution, instead ofturning brown, became white, and no film formed. This was explained by in-creased hydrolysis of the Bi/triethanolamine solution on aging; it seems that somehydrolysis is good for the process, but too much prevents adherence and thereforefilm formation.

A solution similar to the previous ones was used in Ref. 27. One major dif-ference between the films obtained in this study and the previous ones is the film


thickness—ca. 1 �m thick after 30 min at 95°C, compared to an order of magni-tude less after 40 min in a boiling solution, then 4 hr at room temperature. Themost likely difference, in the absence of a knowledge of the triethanolamine con-centration used in the present study, is that this triethanolamine concentration waslower in the present study and therefore thicker films could be obtained in ashorter time, since the the Bi would be less complexed. This is supported by thefact that the Bi(NO3)3 was triturated with the triethanolamine, which suggests thatonly enough triethanolamine was added to dissolve the Bi(NO3)3.

6.4.2 Bi2Se3 (See Table 6.3)

There are only a few papers dealing with Bi2Se3, and therefore they will all betreated in one section. The first report was based on the early Bi2S3 depositions,using a triethanolamine/ammonia bath and selenosulphate [34]. Films were de-posited both with and without hydrazine. The deposition was faster with hy-drazine, as may be expected, although in this case the films lost adherence if leftin the solution more than ca. 30 min. Also, while the films deposited from the hydrazine bath were single phase, traces of elemental Se were found in the hy-drazine-free bath. The large difference in resistivity between films deposited fromthe two baths is interesting, although no reason for this difference was suggested.In contrast to the corresponding deposition of Bi2S3, where the age of the Bi/tri-ethanolamine solution was important (see earlier), the age of the Bi solution is notcritical to obtaining a good film; in fact, thicker films were obtained using evenstrongly aged Bi solutions [26]. This was ascribed to the low temperature used inthe formation of the selenide compared to the sulphide.

Films were also made using N,N-dimethylselenourea as Se source [35].Na2SO3 was added, as usual for selenoureas, to minimize oxidation of the sele-nourea. X-ray diffraction showed only a very broad and ill-defined spectrum ofthe as-deposited film. As for Bi2S3, annealing at a relatively low temperature(200°C) was sufficient to crystallize the film and show well-defined peaks. The

TABLE 6.3 Bismuth Selenide

Bandgap ResistivitySolution (eV) (�-cm) Type Ref.

Na2SeSO3/TEA-NH3/ 1.15* ca. 104–105 n- 26, 3430�C/pH 9.9 � 0.1

� hydrazine 1.03* 5 � 102

Na2SeSO3/NH3/RT 1.42 28DMSeU*/TEA/RT—40�C 1.7–1.4* �107 35Annealed 200�C 1.57–1.07 �0.1 n- dir.


crystal size in the annealed film, estimated from the peak widths, was ca. 12 nm.The properties, in particular the bandgap calculated from the optical spectra, weredependent on the film thickness. Thus the range of bandgaps given in the tablevary from a film 0.09 �m thick (higher bandgap) to 0.15 �m (lower bandgap).This is a very large range for such a small change in thickness. Since the varia-tions are probably due to size quantization, a measure of the XRD peak widths forthe different thicknesses (of minimally annealed films) would be of interest here.The as-deposited films were very photosensitive (light-to-dark ratio up to 70), butthe annealed films were much less so.

6.5 COBALT

In the early work by Beutel and Kutzelnigg on film formation on variousmetals from hot thiosulphate solutions of various metal compounds, one of thosethat resulted in apparent CD was Co [2]. Since no characterization of the films wasgiven and it is possible that the coloration was due to reaction between the thio-sulphate and the substrate metals, this author made a single experiment, mixingCo2� with excess thiosulphate and heating in a glass vial. A black film formed onthe glass (this was also carried out for Ni and Fe; see later). Although no charac-terization was made of this film, it indicates that Co-S was indeed formed in theexperiment of Ref. 2.

CoS was deposited at room temperature from a triethanolamine/ammonia-complexed solution of CoCl2 using thioacetamide as sulphur source [36]. Bothcompositional analysis (CoS1.035) and XRD analysis showed the formation ofCoS. From the optical spectrum, a direct bandgap of 0.62 eV was found. The filmswere p-type with a resistivity of ca. 106 �-cm.

Metallic grey-brown films of Co3S4 were deposited from a CoCl2/NH3 so-lution using thiourea at temperatures between room temperature and 50°C [36a].The film resistivity was approximately 105 �-cm.

CoSe was deposited using a similar composition to that for CoS, except thatselenosulphate was used in place of thioacetamide, NaOH and hydrazine werealso added to the solution, and the deposition was carried out close to 100°C [37].In contrast to CoS, no XRD pattern was observed for the as-deposited films; an-nealing at 280°C crystallized the films to give a defined pattern corresponding toCoSe. A direct bandgap of 0.45 eV was estimated from the optical spectrum. Thefilms were p-type and the resistivity 104 �-cm.

6.6 COPPER

Copper chalcogenides can be readily deposited by CD. There is a strong affinitybetween Cu and S or Se; metallic Cu exposed to elemental S dissolved in a sol-vent (e.g., dimethyl sulphoxide) will quickly turn black due to formation of cop-


per sulphide. This strong affinity between the elements is manifested as a low-sol-ubility product of the various Cu chalcogenides. A complicating factor in this de-position, however, is the large number of different phases and stoichiometries thatcan exist. Four well-defined room-temperature phases of the CuxS system areknown: chalcocite (x � 2); djurleite (x � 1.96); digenite (x � 1.8), and covellite(x � 1). Mixed phases can also occur, of course. Since only CuS gives an identi-fiable XRD pattern in the CD Cu-S films, a knowledge of which phase(s), otherthan CuS, exists is unknown, although the average composition has been mea-sured in many cases. The various phases can, in some cases, be changed from oneto another after deposition. This has been shown for Cu-Se using electrochemicalpolarization [38] and by aging and then reversed by heating [39]. Treatment ofCuxS film with gaseous H2S apparently changes the composition (possibly alsothe phase) toward a more S-rich film [40].

Cu-S shares with PbS and Sb-S the distinction of being the first publishedCD compound [1]. This and (for many decades) subsequent reports involving CDCu-S described decomposition of thiosulphate solutions of Cu salts to give Cu-Sfilms. These (and other, mainly PbS) films were known as lüsterfarben (lustrouscolors) due to the varied interference colors obtained on metal substrates by de-position of PbS or Cu-S (see Sec. 2.1 and Sec. 5.2 for more details of the historyof these lüsterfarben). As for PbS, very little characterization of these deposits wasreported in those early papers apart from their various colors.

If this early work on CD of Cu-S was driven by the attractive colors they im-parted to metallic substrates, more recent studies were driven initially by their po-tential use in Cu2S/CdS photovoltaic cells (these cells are no longer studied to anyextent due to their perceived instability, although, with what has been learnedabout Cu-containing chalcopyrite, such as CuInSe2, thin-film cells over the pastcouple of decades, it would not be surprising if such studies were again pursued)and later in solar control coatings (see Sec. 2.13).

Thiosulphate and sulphite are sufficiently reducing to reduce Cu2� to Cu�.Therefore the Cu in solutions of Cu2� containing sufficient thiosulphate, seleno-sulphate, or sulphite should be predominantly in the monovalent form. This wouldlead to the expectation that the main product will be something close to Cu2S(e).While this is often the case, CuS(e) is deposited in some cases. However, it is ar-guable whether this reduction of Cu2� is, in fact, important in practice. The rea-son is based on an XPS study that showed that Cu in its compounds with S, Se,and Te is normally in the monovalent state; it is the chalcogenide ion (or polyion)that is believed to change oxidation states in these compounds [41].

An interesting characteristic of CD Cu-S films deposited from thiosulphatesolution is the range of compositions that can be obtained by varying the deposi-tion conditions [40]. Elemental analyses of the precipitated CuxS powders ob-tained by heating a solution 0.1 M each in CuSO4 and Na2S2O3 showed that thecomposition varied from x � 1.7 to x � 1.0, with longer reaction times and higher


temperature giving lower values of x. X-ray diffraction of these end-members wasconsistent with a phase where 1.86 � x � 1.96 (therefore the Cu1.7S was a mix-ture of predominantly this phase, presumably djurleite, and other phases) and x �1.0, respectively. XPS of films deposited from the same solutions showed the sul-phur of the CuS to have a higher binding energy (more oxidized) than that of theCu1.7S, in agreement with a variable valence state of the chalcogen in these com-pounds. Variation of the Cu:thiosulphate ratio also affected the composition.Compositions corresponding to Cu2S (Cu:thiosulphate 1:2); Cu1.8S (1:1); Cu1.4S(1:2.5), and CuS (1:3) were obtained, although in film form, only CuS gave anXRD pattern [42,43]. The variation in composition is not a monotonic function ofthe Cu:thiosulphate ratio. For equimolar Cu and thiosulphate, there is not enoughthiosulphate to both reduce all the Cu2� and form the sulphide, hence a partiallyreduced Cu2� is formed; while for a ratio of 1:2, there is enough thiosulphate toboth reduce all the Cu2� to Cu� and to form the sulphide. More difficult to un-derstand at first sight is why an increasing concentration of thiosulphate appar-ently oxidizes the Cu-S to an increasingly greater amount of CuS, until at a ratioof 1:3, only CuS is formed. However, it must be kept in mind that CuS is not sim-ply composed of divalent Cu, but is either 2Cu�S2

2� or a mixed-valence com-pound with both S2� and S2

2� groups. Thus it may be more useful to consider theeffect of the thiosulphate on the sulphur species. A possible hypothesis is that ex-cess thiosulphate results in the formation of elemental S, which can react with S2�

to give the polysulphide ion, S22�, which exists in CuS. More generally, and re-

gardless of the specific mechanism, a higher concentration of S relative to Cu canbe expected to favor more S in the final product.

The various Cu-S and Cu-Se films generally exhibit similar optical spectrafor comparable thicknesses. Figure 6.1 shows some such spectra. The peak at ca.0.6 �m is characteristic of these films, and the drop in transmission at longer

FIG. 6.1 Optical transmission spectra of various CuxS films. (a) Cu2S; (b) Cu1.8S; (c)Cu1.4S; (d) CuS. (Adapted from Ref. 43, with permission from Elsevier Science (USA)).


wavelengths is presumably due to absorption and reflection by free holes(CuxS(e), as for the other Cu-S(e) compounds, is p-type and in most cases highlydegenerate, therefore relatively highly conductive). The CuS (covellite) phasetends to be somewhat more conductive than the other sulphide phases.

CD Cu-S(e) films have been proposed for a number of different potentialapplications. Solar control coatings, where the visible and IR transmission and re-flectivity can be varied, is probably the most studied, e.g., Refs. 44 and 45. Therelatively high conductivity and the partial transmittance in the visible spectrumare useful for transparent conductors [46]. Other possible applications are forCu2� sensor electrodes and electrical contacts for ceramic devices [46].

In the tables for both Cu-S and Cu-Se (Tables 6.4 and 6.5), the column de-noting conductivity type has been deleted (these semiconductors are always p-type), and, in its place, the phase (and/or composition) has been given. In somecases, particularly for the sulphides, where no XRD pattern was seen (except forCuS), no phase (or composition) was proven and therefore no entry is given in thetable. This was not a problem for Cu-Se, since XRD spectra were always clear anddefinitive.

6.6.1 Cu-S (See Table 6.4)

In Ref. 9, no structural or compositional characterization was given; Cu2S was as-sumed. The films were not uniform and did not adhere to the substrates. The highvalues of resistivity are unusual for this material: it may be that 10�3–10�4 wasintended.

TABLE 6.4 Copper Sulphide

Bandgap Resistivity Phase orSolution (eV) (�-cm) composition Ref.

S2O3/RT/pH ca. 2 1.2 104–105* 9S2O3/60�C/pH 0.5 2.4 dir. 10�4 47S2O3/60 or 70�C/pH 2–5 10�2–10�4* Varied 40S2O3/50�C/pH 5 (AcH) 1.7–2.0 �10�3 (CuS) Varied 42, 43

�3 � 10�3 (Cu2S)S2O3-dimethylthiourea 1.55 indir. 8 � 10�3 CuS 48Anneal 300�C 1.55 indir. 8 � 10�4 Cu1.8SAnneal 400�C 1.4 indir. 10�3 Cu1.96STU/NH3/30°C ca. 2.26 ind. 0.25 49TU/TEA-acetate-NH3/ 2.58 dir. and 3 � 10�3 Cu1.86S 50

40�C/pH 9.4 ca. 1.7*CuCl/TU/EDTA/NH2OH/ 1.45 10�3–10�4 CuxS 51

RT—80�C/pH 8.5–11.5 1.83 � x � 1.85TU/TEA-NH3/RT ca. 1.5 indir. 0.004 � 1 52


In Ref. 47, as before, no structural or compositional characterization wasgiven, and the films were classified as CuxS. Average grain size, measured byscanning electron microscopy (SEM), was ca. 70 nm.

In Ref. 40, films were deposited on various polymer substrates that inmost cases required pretreatment with either various organosilanes or poly-(ethyleneimine) in order for good film formation to occur. This technique appearsto constitute a case of adsorption of colloids from the solution rather than forma-tion of Cu-S directly on the substrate. For example, MNH groups on the imine-coated substrate become protonated in the acidic solution and attract the appar-ently negatively-charged Cu-S particles. The composition varied withtemperature and time of deposition (see the previous general discussion). The re-sistivity of the films varied with the deposition conditions and substrate; CuS wasmore conducting than compositions closer to Cu2S.

In Refs. 42 and 43, composition depended on the Cu:thiosulphate ratio (seethe earlier general discussion). Only CuS gave a measurable XRD pattern. CuSwas less transmitting in the IR region than CuxS (x � 1.4–2.0), and for CuS,greater thiosulphate concentration (constant [Cu]) resulted in less transparency inthe IR, although film thickness was fairly constant. This is expected from the dataof Ref. 40, where greater thiosulphate concentration resulted in a deposit closer toCuS in composition and with higher electrical conductivity (therefore less trans-parent in the IR).

A combination of thiosulphate and dimethylthiourea was used in Ref. 48.Using only thiosulphate resulted in slower deposition and, more importantly,poorly adherent films. The resistivity of the as-deposited CuS dropped to 3 � 10�4 �-cm on annealing at 200°C (in N2) without a change in phase or com-position. Annealing at 300 and 400°C resulted in loss of S and phase changes. Thecrystal size was 11 nm (200°C anneal), 19 nm (300°C), and 20 nm (400°C). Pre-sumably no clear XRD pattern was obtained for the as-deposited film of CuS.

In Ref. 49, the composition of the films was given as Cu1.8S. However, noXRD pattern was found and no compositional analysis given, and therefore it isunclear just what the actual composition and phase were.

A potentiometric technique was used in Ref. 50 to measure composition,found to be Cu1.86S. In this deposition, stirring the deposition solution resulted innonuniform and poor-quality films, while good films were obtained in unstirredfilms. The bandgap was measured to be direct, with a value of 2.58 eV. This is aparticularly high value. Examination of the transmittance spectrum showed asharp drop in transmission at ca. 1.7 eV, which is more likely to be the truebandgap.

In Ref. 51, the composition was CuxS, with 1.83 � x � 1.85. The films werethicker than most others—several microns. The terminal thickness was very pHdependent; at pH 8.5, it was ca. 0.5 �m, and at pH 11, an order of magnitude


greater. A pH of 10 was optimum in terms of maximal thickness without exces-sive bulk precipitation.

No structural or compositional data was given in Ref. 52. Although avalue of bandgap was not given, from the optical spectra, an indirect bandgap ofbetween 1.5 and 1.6 eV could be estimated. The specific resistivity dropped withincreasing film thickness by more than two orders of magnitude, � 1 �-cm fora thickness of 0.15 �m to 4 � 10�3 �-cm for 0.35 �m. It was suggested thatthis was due to increasing nonstoichiometry as deposition proceeded. If a sec-ond deposition was carried out on a previous one, the resistivity (measuredacross the film) was characteristic of a single layer, suggesting that a relativelyinsulating layer was deposited on the first layer in the early stages of the seconddeposition.

This group published a number of other papers on CuxS, using the same ba-sic deposition solution, with emphasis on variations of their spectral properties forpossible use in solar control coatings. Some examples are given in Refs. 44, 45,and 53 (the last deals with films deposited on Kapton foil). This last showed someweak CuS (covellite) peaks in the XRD spectrum; no inference of whether thiswas the major phase or not could be made.

6.6.2 Cu-Se (See Table 6.5)

The three main phases encountered in CD Cu-Se films are Berzelianite (Cu2�xSe,where x is typically ca. 0.2); Umangite (often written as Cu3Se2 but may be con-siderably lower in Cu) and Klockmannite (CuSe).

In Ref. 54, XRD showed the deposit to be hexagonal CuSe. Analysis of theabsorption spectrum gave a direct bandgap of 2.02 eV. As commonly seen forthese compounds, there was still strong absorption at lower energies (e.g., at 1.9eV, the absorption coefficient was �7 � 104 cm�2), possibly due to an indirecttransition but likely due, at least in part, to free-carrier absorption. From Hall mea-surements, the doping (acceptor) density was ca. 1022 cm�2 (heavily degenerate)and the mobility ca. 1 cm2V�1sec�1. The dependence of film thickness and depo-sition rate on the deposition parameters has been studied in a separate paper [62].

Nitrilotriacetate was used as complexant in the deposition in Ref. 55. Cu-Secould be both electrodeposited and chemically deposited from this solution. Theelectrodeposited film was Cu1.8Se with the berzelianite structure, while the CDone was Cu1.2Se with the umangite structure. The XRD pattern of the CD filmsshowed sharp peaks (instrument broadening) with no preferential texture. Elec-tron microscopy of these films (Fig. 6.2) shows large (micron scale) particles that,from their faceted shape and together with the sharp XRD peaks, appear to be sin-gle crystals. This is a particularly large crystal size for a CD film; from this and


the other studies on CD Cu-Se, it can be seen that Cu-Se has a tendency to formrelatively large crystals. This may be due, at least in part, to the high mobility ofCu� (although the tendency is much less for Cu-S) and the relatively low meltingpoints of the Cu-Se compounds in general.

FIG. 6.2 Scanning electron micrograph of a Cu1.2Se film deposited from a selenosul-phate solution of Cu2� complexed with nitrilotriacetate.

TABLE 6.5 Copper Selenide

Bandgap ResistivitySolution (eV) (�-cm) Phase Ref.

SeSO3/TEA � NH3/RT ca. 2.0 dir.* 10�3 CuSe 54SeSO3/NTA*/RT/pH � 9 Cu1.2Se umangite 55SeSO3/TEA � NH3/ 1.20 dir. 10�3–2 � 10�2 Cu1.86Se berzelianite 56

�95�C/pH ca. 10SeSO3/NH3/�45�C/ ca. 10 �-sq. Cu2Se 57

pH � 10SeSO3/pH � 10 3 � 10�3–10�4* Cu2Se or CuSe* 58SeSO3/TEA � NH3/75�C Cu2�xSe berzelianite* 38SeSO3/NH3/RT �2.36 dir. 2 � 10�4 Cu2�xSe berzelianite 59

�1.9 indir.*SeSO3/citrate/5–27�C/ See text ca. 2 � 10�3 Cu2�xSe berzelianite 39

pH � 7 (for all films) Cu3Se2 umangiteSeSO3/citrate/60�C

pH 9 2.37 Cu3Se2 umangite 60pH 12 2.0 CuSe klockmannite

DMSeU/tartrate/50�C 10�3–10�4* CuSe 61DMSeU/tartrate/50�C ca. 2.13 �2 � 10�4 CuSe klockmannite 59


The main difference between Ref. 56 and Ref. 54 is the higher deposition tem-perature of the former. Elemental analysis gave a composition of 65% Cu and 35%Se (at.%) corresponding to Cu1.86Se. X-ray diffraction confirmed the cubicberzelianite phase.

The deposition in Ref. 57 appears similar to those of Refs. 54 and 56, exceptthat only ammonia was used to complex the CuSO4 and the deposition tempera-ture was ca. 45°C. The films were deposited on polyester (overhead transparency)films. From XRD, a composition of Cu2Se (no specific phase given) was assigned.The transmission spectrum was similar to those of Cu-S in general, with a strongfree-carrier absorption beginning in the near IR and a strong absorption (presum-ably bandgap) onset at ca. 700 nm.

Reference 58 involved a similar deposition as the previous case, but with-out ammonia. The composition of the Cu-Se was Cu2Se (using a Cu:selenosul-phate ratio of 1:1) and CuSe (1:5). (This is the same trend as found for sulphidesdeposited from thiosulphate solutions.) The transmission spectrum of the Cu2Sewas similar to that in the previous study, while that of the CuSe showed an ab-sorption onset at ca. 600 nm. Both films showed strong apparent free-carrier ab-sorption starting in the near IR (for the CuSe, even at somewhat shorter wave-lengths). However this absorption appeared to be stronger for the Cu2Se than forthe CuSe, although the resistivity of the former (3 � 10�3 �-cm) was higher thanthat of the CuSe (10�4 �-cm).

Copper acetate was used in Ref. 38; it was noted that if chloride was usedinstead of acetate, no deposition occurred, and this was attributed to adsorption ofchloride on the substrate (Pt). The berzelianite phase with a small amount ofumangite impurity was obtained. The composition and phase of the film could bealtered by electrochemical cathodic polarization (in an aqueous K2SO4 solution).Initially, there occurred an increase in lattice parameters and decrease in x(Cu2�xSe). With continued polarization, a phase change occurred until eventuallyonly orthorhombic Cu2�xSe was present in the film. The umangite phase also dis-appeared, and it was believed that this impurity phase catalyzed the phase trans-formation. The change in composition during cathodic polarization was attributedto reduction of zerovalent Se to Se2�, which was dissolved in the solution. Basedon the study of Folmer and Jellinek [41] discussed earlier, this explanation can beinterpreted as reduction of Se2

2� (“monovalent” Se) to Se2� (divalent Se).Reference 59 provides a comprehensive explanation of the optical spectra

and extracted bandgaps. The direct bandgap of ca. 2.36 eV is compared to the lit-erature value of ca. 2.2 eV and explained by size quantization in the fairly small(20 nm) crystals. An indirect bandgap of 1.9 eV was measured (literature value �1.4 eV), but it was stressed that this provided an upper limit only, since the ab-sorption in this region was dominated by free-carrier absorption, which maskedthe indirect absorption. Annealing decreased the conductivity and the free-carrierabsorption and changed the indirect gap to �1.3 eV.


In Ref. 39, citrate was used a complexant and the pH was lower than in othersimilar studies. The depositions were carried out at around room temperature or,if the deposition was too fast, at lower temperatures (no difference in the nature ofthe films was found with different temperatures). The composition and structureof the deposit was found to be pH dependent: At a pH ca. 7 or lower, berzelianiteCu2�xSe was deposited, while at a slightly higher pH (7.8), the product was uman-gite (ca. Cu3Se2). The exact compositions varied with change in ratio between Cuand selenosulphate concentrations. As with the other Cu-Se and Cu-S films,bandgap determination was complicated by the strong free-carrier absorption. Di-rect bandgaps of 2.2 eV (Cu2�xSe) and 2.8 eV (Cu3Se2) were measured from thetransmission spectra. However, from examination of these spectra, it can be in-ferred that a strong absorption, not arising solely from free carriers, occurred atlower energies. An approximate reanalysis of the transmission spectra, taking intoaccount free-carrier absorption, allowed estimation of indirect bandgaps of1.5–1.6 eV (Cu2�xSe) and 2.0–2.1 eV (Cu3Se2).

The berzelianite phase was subsequently found to slowly transform to theumangite one under ambient conditions [63]. By heating at 140°C in air, this phasetransformation could be reversed. These phase changes could be repeated in acyclical manner.

In Ref. 60, the differences in the two solutions giving Cu3Se2 and CuSe werethe lower pH (9) and higher Cu and citrate concentrations (6 mM) for Cu3Se2,compared to pH � 12 and Cu (and citrate) concentrations of 4 mM (constant se-lenosulphate concentration of 30 mM in both cases). The films were deposited onflexible polyester substrates. It was noted that the deposition was paralleled for themost part by bulk precipitation. The average crystal size of both Cu3Se2 and CuSewas 42 nm.

In Ref. 61, N,N-dimethylselenourea was used (together with CuCl2 insteadof CuSO4 commonly used in the selenosulphate depositions). Film (specific) re-sistivity dropped as thickness increased.

Reference 59 is similar to the previous study. Values for film resistivitywere not given, but it was noted that the films were less conductive than Cu2�xSefilms (2 � 10�4 �-cm) made in the same study using selenosulphate instead ofthe selenourea.

6.7 INDIUM

Indium is very readily hydrolyzable, with a pKa of 4.0, and forms the hydroxideeven in moderately acidic solutions (see Sec. 1.1.2). This means that unless depo-sition is carried out in strongly acidic solution, some hydroxide is likely to be pre-sent in any chalcogenide formed by CD. This is indeed the case in most studies ofIn2S3 deposition reported up to now.

Films of In2S3 on glass deposited from a solution of InCl3 and thioac-etamide were described as early as 1976 by the Kitaev group [64]. Broad XRD


peaks corresponding to �-In2S3 with a crystal size of ca. 5 nm were obtained. Abandgap of 2.45 eV could be estimated from the optical spectra. The resistivitywas very high—ca. 1012 �-cm at room temperature—with an activation energyof 1.2 eV (implying a midgap Fermi level and highly intrinsic material). Filmsannealed at 250°C were photosensitive, with a photoconductivity maximum atca. 500 nm (2.5 eV). It was noted that this was blue-shifted by 0.13 eV, com-pared to single-crystal In2S3. The small size of the crystals suggests that sizequantization occurs here.

In2S3, or, as shown later more probably In(OH,S), was deposited by CD foruse as a buffer layer in photovoltaic cells [65,66]. The deposition bath was againInCl3 and thioacetamide operating at a temperature of 70–80°C. Analysis of thesefilms by XPS showed that oxygen was present in the films, presumably as hy-droxide [67]. Importantly, the results were inconsistent with a mixture of sulphideand hydroxide (which might be expected from this bath) and suggested rathersome compound formation. (Details of photovoltaic cells using these films aregiven in Chap. 9.)

In the first of a series of studies, the same basic bath as previously, but us-ing acetic acid to adjust the pH (to ca. 3, probably somewhat lower—see follow-ing reference) was used [68]. The main purpose of the acetic acid is probably tolower the pH and therefore to reduce the In3� hydrolysis, although its (weak) com-plexing ability with In3� might also play some role in minimizing this hydrolysis.The color of the film (and homogeneous precipitate) varied from whitish yellowto yellow (the bandgap of In2S3 is ca. 2.4 eV), and XPS analysis showed that theS:In ratio was always less than that expected for stoichiometric In2S3 (1.5), al-though it was higher when the thioacetamide and acetic acid concentrations wereincreased. Higher thioacetamide concentration increased the concentration of sul-phide formed, while more acetic acid decreased the hydrolysis to hydroxide. Theresulting films were therefore believed to be composed of both sulphide and hy-droxide, designated as Inx(OH)ySz. The bandgaps (indirect) varied between 2.0 eV(highest S content) to 2.5 eV. For the film with the highest S content, the value of2.0 eV is probably an underestimation, particularly taking into account the yellowcolor typical of a film with a bandgap of at least 2.3 eV. The films were highly re-sistive, between 107 and 108 �-cm.

An XPS investigation of these films was carried out [69]. The pH was moreaccurately measured to be between 2.2 and 2.5. Also it was noted that, althoughhigher concentrations of acetic acid minimized the codeposition of hydroxide,above 0.1 M acetic acid, the films were not homogeneous and poorly adherent.From the XPS spectra, it was concluded the films were of the compositionIn(OH)S, with small variations in the S:OH ratio. Sulphate, probably as surfaceoxidized In-S, was also present.

Structural (XRD and microscopic) studies of the films allowed more defi-nite assignments to be made as to the identity of the films [70]. The structure wasdependent on the thioacetamide and acetic acid concentrations. At low concentra-


tions a composition identified as In5S4 was obtained, while at higher concentra-tions a mixture of cubic �- and �-phases resulted. The former possessed a string-like morphology, while the latter was typically composed of more or less spheri-cal granules. The XRD peaks of the In5S4 deposit were broad, equivalent to acrystal size of ca. 5 nm, verified also by TEM. The In3S3 deposits were apparentlyof somewhat larger crystal size (�10 nm).

A study of the species present in these solutions and the mechanism of thedeposition has been presented [71]. Under the conditions of the depositions, themain solution indium species (in the absence of thioacetamide) are In-Cl (mainly[InCl2]�) complex species. Only ca. 1% of the total In content is present as freeIn3�. No In(OH)3 or hydroxy-complexes were calculated to be present if aceticacid was present (in the absence of acetic acid, the hydroxide could form). Froma kinetic analysis of the deposition reaction, it was concluded that the depositionoccurred by direct reaction between the thioacetamide and the chloro-indiumcomplexes. It was noted that thioacetic acid was the main by-product and that noacetamide was detected (see Sec. 3.2.1.3 for a description of the possible mecha-nisms and by-products of thioacetamide hydrolysis). Acetonitrile (CH3CN), a lesscommon by-product, was also detected at the higher pH values (these depositionstook place between a pH of 2 and 3) but not at the lower ones.

A different study, using essentially the same deposition solution (InCl3 �thioacetamide) at a pH of 3.1 or lower, has been described [72]. The films, de-posited on ITO/glass adhered to the ITO side but not well to the glass. In thiscase, compositional analyses showed the films, which gave electron diffractionpatterns and XRD spectra characteristics of In2S3 (mainly the �-phase, possiblytogether with the �-phase) to be slightly S rich (the precipitate formed in solu-tion was more or less stoichiometric). Thus, although there was no evident dif-ference in the deposition parameters, these films appear different than the mixedsulphide-hydroxide ones described previously. Microscopic investigationshowed the films to consist of a mixture of round particles and needles, forminga porous, spongelike morphology. The films exhibited an increase in bandgap,together with decrease in crystal size with decreasing deposition temperature,due to size quantization (see Chap. 10 for more details). Decrease in solution pHalso resulted in a decrease in the bandgap. The bandgap varied between ca. 2.3and 2.7 eV.

6.8 IRON

No clear-cut example of Fe-S has been described. The closest is in the report ofRef. 2, where, among other metal salts, a boiling aqueous solution of iron thiosul-phate imparted coloration to an iron substrate [2]. No confirmation was given asto the composition or nature of this coloration, and a few attempts by the authorto deposit Fe-S on glass by heating thiosulphate solutions of iron salts were un-successful (unlike the corresponding Co and Ni cases).


6.9 MANGANESE

The stoichiometry of Mn-S precipitated from solution is normally MnS. The sta-ble form of MnS is green �-MnS, which has the rocksalt structure. However, thepink form, which is the form that usually precipitates from solution, is a mixtureof �-MnS (zincblende) and �-MnS (wurtzite), both of which are metastable.MnSe behaves analogously.

MnS was deposited from a room-temperature solution of Mn(II) acetatecomplexed with triethanolamine and buffered with NH4Cl [73]. Thioacetamidewas used as a sulphur source, and hydrazine was also used (it was not specifiedwhether the reaction proceeded in its absence). No XRD pattern was seen in theas-deposited (grey-pink) film; annealing at 500°C in an inert atmosphere gave apattern corresponding to MnS. A bandgap (indirect) of 3.25 was measured fromthe optical spectrum. The film was p-type with a resistivity of ca. 105 �-cm.

Optimization of the film growth from the foregoing bath was carried out[74]. In contrast to many other CD reactions, the growth rate decreased slightlywith increasing temperature (also, the terminal thickness was greater at lower tem-perature—a common occurrence due to reduced bulk precipitation). It may be thatbulk precipitation was so rapid at higher temperatures that the thickness of the filmdeposited at the higher temperatures was less by the time of the first thicknessmeasurement (10 min). While, as with the previous study, no XRD pattern wasseen in the as-deposited film on glass, a clear pattern was observed for films de-posited on SnO2-conducting glass, showing a mixture of the cubic �- and hexag-onal �-phases. However, TEM/ED showed the presence of 3- to 4-nm-sized MnScrystals on glass. The optical bandgap (for the film on glass) was estimated to beca. 3.0 eV and direct. There does not seem to be a clear-cut value in the literaturefor the bandgap of MnS, but it is has been given as 3.0 � 0.2 eV. The films showedonly weak photoconductivity.

MnS has also been deposited from an alkaline (pH 9.7–9.8) thiosulphatebath, using MnCl2 [75]. The deposition was carried out at room temperature afterinitial heating at 70°C (this initial heating step was noted to be essential, althoughno explanation for this was given). The XRD spectrum was barely indistinguish-able from the noise; there was a possible correlation with the spectrum of �-MnS.Optical studies showed a direct bandgap of 3.1 eV, and the resistivity was mea-sured to be between 107 and 108 �-cm.

6.10 MOLYBDENUM

There are two main sulphides of Mo. The stable form is the black, layered MoS2,commonly used as a lubricant. Precipitation from (acidic) solution normally givesthe amorphous MoS3, which converts to MoS2 on heating. An important issuewhen using molybdates as a source of Mo is that solutions of molybdates do notprecipitate the sulphide (selenide) when sulphide (selenide)—either as H2S(Se) or


as an alkaline sulphide (selenide)—is reacted with the molybdate solution, butrather form soluble thio(seleno)molybdate ions, such as MoS4

2�.Chemical deposition of both MoS2 and MoSe2 has been reported from am-

monium molybdate solution [76,77]. For the sulphur and selenium sources,thioacetamide and selenosulphate were used, respectively. Ammonium hydroxidewas added to the sulphide solution, while an acetic acid/ammonium acetate bufferwas used with the selenide solution (pH values were not given). Reducing agents(either hydrazine [76] or sodium dithionite [77]) were added to the baths. Deposi-tion was started at 90–100°C, followed by lowering to room temperature.

No XRD pattern was observed for the Mo-S deposit, but after heating (ap-parently in the deposition solution) in an autoclave at 300°C, the XRD pattern ofMoS2 was obtained. The XRD pattern of MoSe2 was obtained for the as-depositedfilm. It is possible that the as-deposited Mo-S was MoS3, which is often obtainedin an amorphous form from solution reactions at relatively low temperatures andconverts to crystalline MoS2 on annealing.

The estimated bandgaps for the two materials were 1.17 eV indirect (MoS2)and 1.14 eV direct (MoSe2). The latter is unusual, since this is the approximatevalue of the indirect gap of MoSe2; the direct gap is substantially higher.

While the role of the reducing agents (hydrazine and dithionite) was not ex-plicitly discussed, it must be assumed that they play an essential role in formingthe Mo chalcogenides rather than the soluble thio(seleno)molybdate ion.

MoSe2 was deposited from a Mo(VI) (the source used was not specified) so-lution complexed with ammonia to give a hexammine complex and mixed with hy-drazine and selenosulphate at 40°C [78]. The as-deposited films were XRD amor-phous but converted to crystalline MoSe2 after annealing in N2 at 380°C. Elementalanalysis showed the as-deposited films to be nearly stoichiometric MoSe2. A directbandgap of 1.48 eV (1.36 after annealing) was measured. The films were n-typewith a resistivity of ca. 4 � 103 �-cm (ca. 1 �-cm after annealing).

6.11 NICKEL

Ni-S behaves rather similarly to Co-S (see Sec. 6.5, Cobalt). Note that the freshlyprecipitated monosulphides of both metals transform in solution to a more insol-uble form–possibly M(OH)S.

As for cobalt (see earlier), an early study of the coloration of metals by im-mersion in boiling metal salt–thiosulphate solutions resulted in coloration of themetals [2]. Also as for cobalt, a single experiment by the author repeating (simi-lar, not necessarily identical conditions) this experiment, only on glass, rather thanon metals which might be colored by the thiosulphate alone, resulted in a blackfilm on the glass.

Other than the foregoing, only one other paper was found dealing with NiS(and also NiSe) [79]. The baths were based on NiSO4, triethanolamine, and am-


monia. For the sulphide, thioacetamide was used, while for the selenide, seleno-sulphate was the Se source, and NaOH and hydrazine were added to the bath. De-positions were carried out at room temperature. XRD confirmed the formation ofNiS and NiSe films. The bandgaps (direct) were 0.35 eV (NiS) and ) 0.23 eV(NiSe). The films were both p-type, with resistivities of 10 �-cm (NiS) and 0.1�-cm (NiSe).

6.12 SILVER

6.12.1 Ag2S (see Table 6.6)

The Ag� ion forms strong complexes with thiourea (log K � 12.7 for theAg(thiourea)3

� complex and with thiosulphate (log K � 13). The strong bindingof Ag to thiosulphate is exploited in the use of Na2S2O3 solution to remove ex-cess Ag during the developing of photographic films. For this reason, Ag� can

TABLE 6.6 Silver Sulphide


S2O3/RT/pH 2–3 1.2 106–107 9S2O3/RT/pH 2.6 (2.2)* 0.95 dir. 103–104 80� EDTA 0.73 dir. 104–105

S2O3/NH3/50�C/pH 9–11 �2.2 �10 11TU/S2O3

2�/RT/pH 10.1* 2.3 dir. 10�2 p- 81TU/8–55�C/pH 8–10 (NH3) — 103–105* n- 82, 83� EDTA ca. 0.8 dir. 2 � 104

TU/NH3, kinetic study 84TU/NH3, thermodynamic 85

analysisTU/NH3/NH4

�/pH ca. 11, 86structural study

TU/40–80�C/pH 9, 87study of mechanism

TU/40–80�C/pH 9 ca. 1.0* 88TU (NH4OH) Hg2� doped 0.8 (from ca. 103 (dark) n- 89, 90

photocond.)TA/RT/8–55�C 0.95 dir. ca. 105 n- 91� EDTA ca. 104

Dip technique Ag/S2O3/RT 0.83 indir. (100 �) n- 92� TU/80�C/pH 8–11*


be kept in solution under alkaline conditions without the need for another com-plexant. The strong binding of Ag to sulphur in thiourea and thiosulphate sug-gests that the mechanism of Ag2S formation may be of the complex-decompo-sition type rather than through formation of free sulphide. Thus any bondbreaking involving the Ag-complex is intuitively expected to occur at the SMCor SMS bonds of the thiourea or thiosulphate, respectively (see Sec. 3.3.3.1,Eqn. 3.5.5 and discussion following for more details on this topic).

Two forms of Ag2S exist—�-Ag2S (acanthite), a monoclinic form, and �-Ag2S (argentite), which is cubic. �-Ag2S is the form that is stable at room tem-perature and is invariably the one that occurs in CD films.

Thiosulphate was used in an acid bath in Ref. 9. This study covered manysulphides deposited using thiosulphate, and little detail was given on the deposi-tion or on the films themselves.

Equimolar quantities of Ag� and thiosulphate were used in Ref. 80 (as forthe previous study), so the complexation of the Ag� by thiosulphate was not asstrong as it would have been in an excess of thiosulphate. A suggested mechanismfor the deposition was reduction of elemental S to sulphide, formed in the acidicthiosulphate solution, by the moderately reducing thiosulphate. It was stressedthat the thiosulphate was slowly added to the AgNO3 solution with heavy stirring,with the implication that otherwise the thiosulphate would be oxidized. The filmswere reported to be rough rather than the smooth specular films often obtained byCD. Differences in properties were obtained if EDTA was added as an additionalcomplexant. The films with EDTA were somewhat thinner (0.14 instead of 0.19�m), and XRD of the EDTA-free films gave sharp peaks, while those depositedfrom an EDTA-containing bath apparently showed no pattern, hence were proba-bly either very small crystalline or amorphous. As seen from Table 6.6, both ap-parent bandgap and resistivity changed on addition of EDTA to the bath. Thehigher resistivity with EDTA was explained by the smaller crystal size. The lowerbandgap is less obvious; very small crystal size would increase the bandgap, andan amorphous semiconductor has often (although not always) a higher bandgapthan the crystalline form.

The deposition in Ref. 11, from an alkaline thiosulphate bath, was reportedin the context of a general description of deposition of various materials by CD,and only a little characterization was reported. X-ray diffraction showed someAg2S peaks. Optical spectroscopy showed a gradual decrease in transmission overa wide spectral range, and it would be difficult to extract a reliable value for thebandgap from the spectrum.

Note that the deposition in Ref. 81 used thiosulphate as a complexing agentand not ostensibly as a source of S. The thiourea concentration is critical. Thioureais added to the Ag�/S2O3 just until some solid Ag2O is formed. Too little thiourearesults in thin, scattering films. Too much results in films that in their as-depositedstate are good, but after annealing (300°C), voids form. The pH is also critical:


Under 10 the deposition is slow, while above 10.2 it is too fast and bulk precipi-tation dominates.

In Refs. 82 and 83, this deposition was employed both with and withoutEDTA, a strong complexant. The deposition was studied at various depositiontemperatures. Better adhesion was obtained at low temperature (8°C). The resis-tivity of the films was dependent on deposition temperature: 2.103 �-cm (8°C)and 1.5 � 105 �-cm at 25°C. For a film deposited from an EDTA-containing bathat 8°C, it was ca. 2 � 104 �-cm.

The reaction kinetics were analyzed in an early study [84]. Of particularnote is the unusually high activation energy (160 kJ/mole)—about twice the nor-mal value for reaction-controlled CD processes. This contrasts with the muchsmaller value (20.4 kJ/mole) measured in Ref. 87, although there were several dif-ferences in the deposition: The present solution contained ammonia, the pH washigher—probably between 11 and 12 (compared to the borate buffered solutionwith pH � 9 of Ref. 87), and the activation energies were measured at lower tem-peratures. It is interesting that the pH could not be measured directly in this studyusing a pH meter since the pH electrode was apparently rapidly coated with Ag2S.Based on the kinetic study, the overall reaction proposed was

2Ag(N2H4CS)3� � 2OH� → Ag2S � 5N2H4CS � NH2CN � 2H2O (6.1)

While both ammonia and thiourea were present, it is probable that with sufficientthiourea present, the main complexant was thiourea (see the next paragraph). Thiswould also explain the observed dependence of deposition rate on thiourea con-centration: Initially the deposition increased (since thiourea is the source of S),reached a maximum, and then decreased (due to increasing complexation).

Reference 85 presents the thermodynamic side of the previous paper. It ispointed out that although both ammonia and thiourea are present in the solution,because of the much higher stability constant of the Ag-thiourea complexes com-pared to the Ag-ammines, essentially all the Ag will be present as a thiourea com-plex. In this case, it can be assumed that the role of ammonia is only to control pH.

An interesting observation in Ref. 86 was that the density of nuclei formedin the early stages of film deposition did not change with time. The film developedby growth of the initially relatively small (ca. 20 nm) nuclei. This suggests an ion-by-ion type of growth rather than a cluster one.

Reference 87 is a mechanistic study of Ag2S deposition from a thiourea bath(buffered to pH 9 with a borate buffer). There are some unusual properties of thisdeposition. One, the unusually strong effect of stirring on the deposition rate, hasalready been dealt with (Sec. 3.7) and, together with the measured activation en-ergy of 20.4 kJ/mole, suggests a rate-determining diffusion step in the deposition.Another observation is that the films reach a maximum thickness (between 0.5 and1.6 �m) after 30 min of deposition (the thickness increases with temperature butthe 30 min is, surprisingly, temperature independent) and then become thinner


with further time in the bath, presumably because of loss of adhesion of the filmdue to increasing stresses in the film. The pH is fairly critical: �8.8 slows thegrowth down greatly, while �9.4 results in a fast homogeneous precipitation insolution (such precipitation occurs, even at the optimum pH of 9, in parallel withfilm formation).

The properties of the foregoing films (only without the borate buffer) weredescribed in Ref. 88. The XRD peaks of the as-deposited films were narrow andsharp, evidencing relatively large crystal size. The bandgap of the as-depositedfilm, measured from transmission spectra, was ca. 1.0 eV but varied somewhat withdeposition temperature: At 40°C it was 0.91 eV and reached a maximum of 1.02 at60°C and then decreased slightly at still higher temperatures. The absorption onsetwas sharper at higher temperatures, which was interpreted as being due to denserfilms. The refractive index was also slightly higher for higher deposition tempera-ture, again explainable by the same rationale. The films, annealed in N2 at 250°C,were photoconducting, and the photoconductivity spectrum was similar to the ab-sorption spectrum. Time-resolved microwave conductivity measurements werecarried out on the films. Fast decay times and moderately good mobilities werefound from these measurements. In particular, the mobility was very temperaturedependent, and the highest value (5.3 cm2V�1sec�1) was obtained for annealedfilms that had been deposited at 60°C; both higher and lower deposition tempera-tures gave much lower mobilities. While not understood, this dependence empha-sizes the need to optimize these films specifically for any particular application.

A study of the photoconductivity of Ag2S doped with Hg2� or Au3� was de-scribed in Ref. 89. Illumination decreased the resistivity typically 2–3 times. Pho-toconductivity spectra showed best results for Hg doping; Au doping gave ahigher peak sensitivity but a narrower spectrum, with lower sensitivity at shorterwavelengths compared even to undoped films [90].

Reference 91 involved an acid bath (although pH was not given) usingthioacetamide as the source of S. The terminal film thickness was greater forlower-temperature deposition; films thicker than 3 �m were obtained at 8°C. Incontrast to the acid thiosulphate bath, the use of EDTA decreased the resistivity,as did deposition at lower temperatures. Photoelectrochemical activity was foundfor these films (see Chap. 9).

Reference 92 describes not a normal CD process, but one closer to theSILAR technique described in Sec. 2.11.1. However, while the SILAR method in-volves dipping the substrate in a solution of one ion (e.g., sulphide), rinsing to re-move all but (ideally) a monolayer of adsorbed ions and then dipping in a solutionof the other ion (e.g., Ag�), the present technique omits the intermediate rinsingstep. This means that a relatively large amount of solution can remain on the sub-strate between dips, and layer formation proceeds much more rapidly than forSILAR, albeit with less control. A typical rate was 4 nm/dip cycle. In this case, avisible layer of Ag2S formed after several dips. Since interference colors were ob-


tained, the films were smooth, because such colors are not seen on rough, highlyscattering films. The film thickness/number of dips increased with increasing pH;however, the best films were obtained at pH � 9.

6.12.2 Ag2Se

Ag2Se films were first deposited from a bath using selenosemicarbazide as a Sesource and thiourea to complex the Ag at 20°C [93]. The films were specular andhad a resistivity of between 2 and 20 �-cm.

The only other true CD of Ag2Se describes films deposited on polyester sub-strates from an ammoniacal AgNO3 solution with selenosulphate at 0°C and a pHof 10–11 [94]. They were strongly (111) textured, with a crystal size of 9 nm. Theoptical bandgap (direct) was estimated to be 1.8 eV, compared to the normal valueof ca. 1.3 eV. This was attributed to size quantization. The absorption spectrashowed considerable absorption (scattering?) at longer wavelengths, which couldbe due to a lower, indirect bandgap if not to scattering. The resistivity of the filmswas ca. 2 � 10�3 �-cm (200 �-sq).

A dip technique in which metallic Ag films were converted into Ag2Se wasdescribed [95]. The Ag film was made by successive dipping of glass substrates ina AgNO3 solution, followed by dipping in a solution of formaldehyde, and was con-verted to the sulphide by treatment with a solution of SeO2. The films were roughand apparently poorly adherent. The resistivity of the films was ca. 103 �-cm.

6.13 THALLIUM

The few cases reported for CD sulphides and selenides of Tl all reported themonosulphide (selenide)—TlS or TlSe. Tl can be monovalent or trivalent, andthese apparently divalent compounds are believed to be mixed-valence com-pounds, with both Tl(I) and Tl(III) present.

6.13.1 TlS

TlS was deposited from a solution of TlNO3, ammonia, and thiourea at room tem-perature (26°C) [96]. X-ray diffraction showed the formation of TlS. Optical spec-troscopy (both transmission and diffuse reflection) allowed an approximatebandgap of 1.0 eV to be estimated. The films were p-type, with resistivity of ca. 2 � 103 �-cm. Photoconductivity was measured (although not quantified) with apeak at ca. 1.2 eV (ca. 1 �m) and extending from 1.0 eV to beyond 1.4 eV.

6.13.2 TlSe

TlSe was first deposited from a solution of thallium(I) acetate and selenosulphatewith added NaOH and hydrazine at room temperature [97]. The initial films were


mirror-like but became thick (ca. 3–5 �m) and matte black with increasing depo-sition. X-ray diffraction confirmed the film to be tetragonal TlSe. The main pur-pose for making these films was to study their photoconductivity. The films werep-type, with a resistivity of 500 �-cm, which decreased by a factor of 2 after an-nealing (150°C in air). Indium doping (by adding In3� to the deposition bath) in-creased the resistivity (of the annealed film) by an order of magnitude, probablyby introduction of compensating donors. The films, particularly the annealed,doped ones, were highly photoconducting, with a maximum photosensitivity(change in conductivity on illumination/dark conductivity) of �107 at a wave-length of 1.1 �m. The response extended to ca. 1.5 �m (0.83 eV) (low-energyside) and ca. 0.7 �m (high-energy side).

TlSe was also deposited from a solution of Tl2SO4 complexed with tri-ethanolamine and ammonia and selenosulphate at 30°C [49]. Tetragonal TlSe wasidentified by XRD. The bandgap was estimated at 1.12 eV; however, the absorp-tion spectrum appears to show two transitions—one (possibly indirect) at �0.9 eVand another at �1.3 eV. The films were p-type, with a resistivity of 105 �-cm.Considering the high carrier concentration measured (almost 1020 cm�3), this re-sistivity value appears unusually high.

6.14 TIN

There are three sulphides of tin: SnS (grey, metallic; usually nonstoichiometric),Sn2S3 (black), and the layered, yellow SnS2. SnS and SnS2 are formed when hy-drogen sulphide is passed into solutions of Sn(II) and Sn(IV), respectively. Theanalogous selenides also exist, although the existence of Sn2Se3 is apparently insome doubt.

Tin forms soluble thio(seleno)anions. The sulphides tend to be soluble invery alkaline solutions.

6.14.1 Sn-S (See Table 6.7)

Deposition of Sn-S (from a thiosulphate bath) was claimed as far back as 1870 [98].This was based on deposition from a boiling solution of “zinnsalz” [probably tinchloride, but not clear whether Sn(II) or Sn(IV)] complexed with tartrate and us-ing thiosulphate as the source of sulphur. Unfortunately, the substrate was brass;since brass will slowly convert to a dark-colored sulphide upon immersion in boil-ing thiosulphate solution and no characterization of the film was made other thanits color (which would vary initially according to thickness, due to interference ef-fects), there is no evidence that the films were really a sulphide of Sn. In fact, an at-tempt to reproduce these results concluded that there was no Sn in the layer [2]. Onthe other hand, SnS2 films were later obtained using a thiosulphate bath, althoughthe solution composition was different (see later discussion of Ref. 99.)


Reference 100 described a different technique than usual, in that it used el-emental S dissolved in a carboxylic acid. Propionic acid was most often used, al-though other carboxylic acids could also be employed. It was noted that best re-sults were obtained when the SnCl2 was added as a powder to a freshly preparedS solution; aged solutions reacted much more slowly, if at all. It was surmised thatthis was due to changes in the nature of the dissolved Sn2� with time, such as tinoxide, or propionate formation and loss of HCl.

The stoichiometry of the deposit was dependent on the water content of thesolution and on the presence of a complexant. In anhydrous solution and withoutcomplexant, a uniform, brown film with the approximate composition Sn2S3 wasformed. Addition of ca. 1% of water resulted in a uniform, slate-gray film ofSn1�xS. This was explained by the increased ionization of Sn species in water andtherefore increase in the concentration of Sn2�. A substantial increase in the wa-ter concentration resulted in patchy films of an irreproducible nature, with bothbrown and gray regions forming. Complexation of the Sn reduced the reaction rateand allowed more time for the growing films to react with dissolved S, resultingin films of approximate composition SnS2 (yellow) together with some Sn2S3. Inthis case, some water was needed to form good films. An interesting and unusualcharacteristic of this deposition method is that if the deposition is allowed to pro-ceed for a long time, the amount of precipitate formed in the reaction is decreasedand the films grow thicker. This was explained through dissolution and reprecip-itation of Sn-S solid phases via soluble thiostannate species. Dissolution was as-sumed to occur preferentially in the bulk precipitate due to its greater accessiblesurface area compared to that of the film.

Another unique characteristic of this process is that a band of Sn-S abovethe level of the deposition solution was frequently observed. This above-solutionfilm was ascribed to reaction of volatile SnCl4 and H2S. The presence of the for-

TABLE 6.7 Tin Sulphide

Bandgap ResistivitySolution (eV) (�-cm) Composition Ref.

S in RCOOH/SnCl2/�90�* See text SnS (p-) 100Sn2S3; SnS2*

S2O3/SnCl4/RT/pH 1–2 2.35 103–104 (n-) SnS2 99TA/SnCl2/AcH/TEA � NH3/RT 1.51 indir. 2.5 � 1010 (n-) SnS (anneal) 101TA/SnCl2·2H2O in acetone 1.3* 2 � 104–107 (p-) SnS 102

TEA � NH3/RT—75�CTA/SnCl4/EDTA � NH3 2.3 dir. 1.2 (n-) SnS2* 103

hydrazine/RT/pH ca. 10


mer could be explained by reactions of the type

2SnCl2 � S → SnCl4 � SnS (6.2)

while the H2S (which could be detected separately) might be formed by SnCl2-re-duction of S.

X-ray diffraction showed the gray (low-water-concentration) films to havethe pattern of SnS, while compositional analysis showed them to be Sn deficientby typically 10%. Crystals were large (micron-sized or larger). The other filmstended to be mixtures of approximate compositions Sn2S3 and SnS2, with a nee-dle-like morphology (typically 1 �m long by 0.1 �m wide).

The SnS had an indirect bandgap of 1.0–1.3 eV and was p-type. It was moredifficult to estimate the bandgaps of the other films due to their mixed nature.However, approximate bandgaps of 1.8 eV (Sn2S3) and 2.4 eV (SnS2) could be es-timated from the optical spectra.

The possible mechanisms of this unique deposition are not considered inChapter 3 and therefore will be done so here. The reaction of metal salts with el-emental S in nonaqueous solvents in which S dissolves is known, even if themechanism is not clear. In the present case, two plausible mechanisms can begiven. The Sn(II)/Sn(IV) redox potential is relatively negative (�0.15 V vs. SHEin aqueous solution; much more negative in alkaline solution, although this is notrelevant in the acidic conditions used here). Since only Sn(II) was added to the so-lution, any Sn(IV) present will be formed in the solution and is likely to occur inlow concentration under most conditions. This means that, from the Nernst equa-tion (see Chap. 1), the potential of the solution will be more negative than the stan-dard potential, possibly by a large amount. Another, more negative potential thatmay be relevant is the Sn2�/Sn0 (�0.14 V; see later). The S/S2� standard poten-tial is �0.45 V. Since only the oxidized form of this couple was added to the so-lution, the redox potential (which for our purposes means the potential wheresome appreciable concentration of S2� will be formed) will be considerably pos-itive of this value. Add to this the known non-Nernstian behavior of the S/S2� cou-ple when [S] � [S2�], when the potential shifts strongly positive to a greater ex-tent than expected from the Nernst equation [104], and it is feasible that the Sn(II)may reduce S to S2� in sufficient concentration to form Sn-S.

By esentially the same reasoning, it could be argued that the Sn(II) might re-duce itself (disproportionate) into elemental Sn (at low concentration). As justnoted, the standard potential of Sn(II)/Sn(0) is �0.14 V. It might be argued thatthermodynamically, this is more likely than the reduction of S (although the non-Nernstian behavior of S/S2� will at least reduce this difference). Metals immersedin nonaqueous solutions of S can react to form a layer of the metal sulphide (therate depending on the metal and on the temperature—e.g., Cu will readily sulphideat room temperature, while sulphidization of Zn will proceed slowly even at hightemperature). The small Sn0 nucleii that may be formed in the disproportionation


of Sn(II) would be chemically very active and more likely to react with the S insolution.

In connection with this mechanism, it has been reported that the reaction be-tween elemental S and Sn in alkylammonium compounds or amines under hy-drothermal conditions gives various organic tin-suphide species and even, undersome conditions (low pH), SnS2 [105]. A possible mechanism for this process wasproposed based on nucleophilic attack of the basic amine or hydroxide on the S8

chain:

S8 � RNH2 → �SMS�SMN�

H2R (6.3)

resulting in formation of polysulphide ions. This mechanism is unlikely in theacidic solutions used for the CD process. However, according to the principles justdiscussed based on the Nernst equation, elemental S in solution may be expectedto contain a (very low) equilibrium concentration of (poly)sulphide. If the sul-phide is removed by reaction (with Sn2�, in this case), then even this very lowconcentration may be enough to sustain the formation of a Sn-S solid phase.Clearly, the lower the solubility product of the metal sulphide, the more likely thisprocess is to occur.

Considering that homogeneous precipitation of metal chalcogenides(mainly sulphides) by reaction between metal ions and dissolved chalcogen is wellestablished, the main difference between this deposition and similar reactionsseems to be that the products adhere to a substrate to give a visible film (in thiscase) rather than only precipitate. Whether this is connected with the redissolu-tion/redeposition process that occurs with the Sn-S system or has some other explanation is important. If the former, it may be limited to only those systems thatbehave similarly. Otherwise it is not unreasonable to expect that other metal sul-phides and selenides (possibly also tellurides, although tellurium tends to be muchless soluble, if at all, in such solvents) may be deposited as films in this manner.

In Ref. 99, yellow-gold films were obtained that gave no XRD pattern, butthe chemical composition (as well as color) was consistent with SnS2. Increase inpH (more than 2) reduced the deposition rate, while increase in temperature led toprecipitation in solution and therefore thinner films.

In Ref. 101, the films in the deposition were deep brown. No XRD patternwas observed, but after annealing in an inert atmosphere at 410°C the pattern ofSnS was obtained and the stoichiometry confirmed by elemental analysis. Thebandgap (1.51 eV, indirect transition) was higher than the literature value (1.3eV), and this was rationalized as resulting from the apparent amorphous structure.The room-temperature conductivity (4.10�9 S-cm�1 � a resistivity of 2.5 � 108

�-cm) is low for a relatively low-bandgap material, suggesting either a very stoi-chiometric and intrinsic material or high grain-boundary resistance. The filmswere photoconductive (the sensitivity was not given), with a spectral range from550 to 1050 nm and a peak at ca. 850 nm.


In Ref. 102, the solution should be turbid; if too much triethanolamine wasadded and the solution was clear, no deposition occurred. Too much NH3 (pH notgiven) led to incorporation of hydroxide. The bandgap was given as ca. 1.3 eV;however, from the optical spectra, there is a variation in the bandgaps, dependingon the deposition conditions and film thickness. In particular, the lower the tem-perature of deposition, the larger the bandgap (which analyses of the spectra showto be indirect and to vary from 1.45 to 1.2 eV). This behavior is typical of sizequantization. The resistivity varied strongly with film thickness; the values shownin Table 6.7 are for 0.35 �m (107 �-cm) and 1.2 �m (2 � 104 �-cm). The filmswere mildly photoconductive (maximum sensitivity � 10).

These SnS films (in one study, propylene glycol was used instead of acetoneto dissolve the SnCl2), coated with CD CuxS, were shown to possess spectral char-acteristics favorable for various solar control purposes [106,107]. Depending onthe thicknesses of the two layers, films with varying absorption and reflection couldbe obtained that might be suitable for solar collectors or for window glazing.

In Ref. 103, the pH of the bath was rather critical; films deposited at pH �10.5 were powdery and poorly adhering, while at pH � 9.5, no deposition was ob-served. No structural or compositional characterization was given, but from thetransmission spectrum it could be assumed that the film was SnS2 (and also highlyscattering).

Films designated as Sn(O,S) were deposited from a solution containingtin(IV) acetate, HCl, and thioacetamide (the concentration of the latter two com-ponents determining the S:O ratio) [108]. These films were prepared as buffer lay-ers for photovoltaic cells (see Chap. 9), and little characterization of the filmsthemselves, other than some XPS, was reported. The XPS results suggested thatthe films were a mixture of SnO2 and some Sn-S species.

While not strictly CD, SnS has been deposited by an immersion techniquewhereby a glass substrate was immersed in a cold sulphide solution, followed,without rinsing, by immersion in a hot SnCl2 solution, and this cycle was repeatedto increase the film thickness [109]. The film properties, in particular the electri-cal resistivity, were very dependent on the pH of the SnCl2 solution.

6.14.2 Sn-SeOnly one example of Sn-Se has been reported [110]. Films were deposited from aroom-temperature selenosulphate solution of SnCl2 complexed with tri-ethanolamine and added NaOH. Polyvinylpyrollidone (PVP) was also added andin general slowed down the deposition. At an optimum concentration of PVP, amaximum terminal thickness was obtained (although no comparison with filmsdeposited from PVP-free solutions was given). No XRD pattern was observed forthe as-deposited films; heating in an inert atmophere at ca. 330°C gave the patternof SnSe. The bandgap was 0.95 eV (indirect). The films were n-type, with a re-sistivity of ca. 10 �-cm


REFERENCES1. C Puscher. Dingl. J. 190:421, 1869.2. E Beutel, A Kutzelnigg. Monats. 58:295, 1931.3. KC Mandal, A Mondal. J. Phys. Chem. Solids 51:1339, 1990.4. O Savadogo, KC Mandal. J. Electrochem. Soc. 139:L16, 1992.5. O Savadogo, KC Mandal. Sol. Energy Mater. Sol. Cells 26:117, 1992.6. JD Desai, CD Lokhande. Thin Solid Films 237:29, 1994.7. RS Mane, BR Sankapal, CD Lokhande. Thin Solid Films 353:29, 1999.8. LP Deshmukh, SG Holikatti, BP Rane, BM More, PP Hankare. J. Electrochem. Soc.

141:1779, 1994.9. CD Lokhande. Mater. Chem. Phys. 28:145, 1991.

10. I Grozdanov, M Ristov, G Sinadinovski, M Mitreski. J. Noncryst. Solids 175:77,1994.

11. I Grozdanov. Semicond. Sci. Tech. 9:1234, 1994.12. MTS Nair, Y Pena, J Campos, VM García, PK Nair. J. Electrochem. Soc. 145:2113,

1998.13. P Pramanik, RN Bhattacharya. J. Solid State Chem. 44:425, 1982.14. RN Bhattacharya, P Pramanik. J. Electrochem. Soc. 129:1642, 1982.15. RN Bhattacharya, P Pramanik. Sol. Energy Mater. 6:317, 1982.16. JD Desai, CD Lokhande. Thin Solid Films 249:135, 1994.17. JD Desai, CD Lokhande. Ind. J. Pure Appl. Phys. 31:152, 1993.18. CD Lokhande, AU Ubale, PS Patil. Thin Solid Films 302:1, 1997.19. JD Desai, CD Lokhande. Mater. Chem. Phys. 34:313, 1993.20. RS Mane, BR Sankapal, CD Lokhande. Thin Solid Films 359:136, 2000.21. S Biswas, A Mondal, D Mukherjee, P Pramanik. J. Electrochem. Soc. 133:48, 1986.22. MTS Nair, PK Nair. Semicond. Sci. Tech. 5:1225, 1990.23. PK Nair, MTS Nair, O Gomezdaza, RA Zingaro. J. Electrochem. Soc. 140:1085,

1993.24. JD Desai, CD Lokhande. Mater. Chem. Phys. 41:98, 1995.25. P Pramanik, RN Bhattacharya. J. Electrochem. Soc. 127:2087, 1980.26. RN Bhattacharya, P Pramanik. J. Electrochem. Soc. 129:332, 1982.27. LP Deshmukh, KV Zipre, AB Palwe, BP Rane, PP Hankare, AH Manikshete. Sol.

Energy Mater. Sol. Cells 28:249, 1992.28. RK Nkum, AA Adimado, H Totoe. Mater. Sci. Eng. B. 55:102, 1998.29. D Mukherjee, AK Dutta. J. Mater. Sci. Lett. 8:511, 1989.30. PK Nair, J Campos, A Sanchez, L Banos, MTS Nair. Semicond. Sci. Technol. 6:393,

1991.31. ME Rincon, PK Nair. J. Phys. Chem. Solids 57:1937, 1996.32. ME Rincon, R Suárez, PK Nair. J. Phys. Chem. Solids 57:1947, 1996.33. PK Nair, MTS Nair. Semicond. Sci. Technol. 7:239, 1992.34. P Pramanik, RN Bhattacharya, A Mondal. J. Electrochem. Soc. 127:1857, 1980.35. VM García, MTS Nair, PK Nair, RA Zingaro. Semicond. Sci. Technol. 12:645,

1997.36. PK Basu, P Pramanik. J. Mater. Sci. Lett. 5:1216, 1986.36a. FC Eze, CE Okeke. Mater. Chem. Phys. 47:31, 1997.


37. P Pramanik, S Bhattacharya, PK Basu. Thin Solid Films 149:L81, 1987.38. C Lévy-Clément, M NeumannSpallart, SK Haram, KSV Santhanam. Thin Solid

Films 302:12, 1997.39. M Lakshmi, K Bindu, S Bini, KP Vijayakumar, CS Kartha, T Abe, Y Kashiwaba.

Thin Solid Films 370:89, 2000.40. T Yamamoto, K Tanaka, E Kubota, K Osakada. Chem. Mat. 5:1352, 1993.41. JCW Folmer, F Jellinek. J. Less Common Metals 76:153, 1980.42. I Grozdanov, CK Barlingay, SK Dev. Thin Solid Films 250:67, 1994.43. I Grozdanov, M Najdoski. J. Solid State Chem. 114:469, 1995.44. MTS Nair, PK Nair. Semicond. Sci. Technol. 4:599, 1989.45. PK Nair, VM Garcia, AM Fernandez, HS Ruiz, MTS Nair. J. Phys. D:Appl. Phys.

24:441, 1991.46. I Grozdanov, CK Barlingay, SK Dev. Mater. Lett. 23:181, 1995.47. KM Gadave, CD Lokhande. Thin Solid Films 229:1, 1993.48. MTS Nair, L Guerrero, PK Nair. Semicond. Sci. Technol. 13:1164, 1998.49. RN Bhattacharya, P Pramanik. Bull. Mater. Sci. 3:403, 1981.50. E Fatas, T Garcia, C Montemayor, A Medina, E Garcia Camarero, F Arjona. Mater.

Chem. Phys. 12:121, 1985.51. AJ Varkey. Sol. Energy Mater. 19:415, 1989.52. MTS Nair, PK Nair. Semicond. Sci. Technol. 4:191, 1989.53. J Cardoso, O GomezDaza, L Ixtlilco, MTS Nair, PK Nair. Semicond. Sci. Technol.

16:123, 2001.54. A Mondal, P Pramanik. J. Solid State Chem. 47:81, 1983.55. G Hodes, D Cahen. Ternary adamantine materials for low-cost solar cells (Project

No. IL-2-04132-1), Quarterly Status Report to Solar Energy Research Institute,Golden, CO, 1985.

56. GK Padam. Thin Solid Films 150:L89, 1987.57. I Grozdanov. Chem. Lett. 3:551, 1994.58. I Grozdanov, CK Barlingay, SK Dey. Integrated Ferroelectrics 6:205, 1995.59. VM Garcia, PK Nair, MTS Nair. J. Cryst. Growth 203:113, 1999.60. B Pejova, I Grozdanov. J. Solid State Chem. 158:49, 2001.61. CA Estrada, PK Nair, MTS Nair, RA Zingaro, EA Meyers. J. Electrochem. Soc.

141:802, 1994.62. A Mondal, P Pramanik. J. Solid State Chem. 55:116, 1984.63. M Lakshmi, K Bindu, S Bini, KP Vijayakumar, CS Kartha, T Abe, Y Kashiwaba.

Thin Solid Films 386:127, 2001.64. GA Kitaev, VI Dvoinin, AV Ust’yantseva, MN Belyaeva, LG Skornyakov. Inorg.

Mater. 12:1448, 1976.65. KO Velthaus, J Kessler, M Ruckh, D Hariskos, D Schmid, HW Schock. In: 11th

ECPV Solar Energy Conf., Montreux, Switzerland, 1992, p 842.66. D Braunger, D Hariskos, T Walter, HW Schock. Sol. Energy Mater. Sol. Cells

40:97, 1996.67. D Hariskos, M Ruckh, U Ruhle, T Walter, HW Schock, J Hedstrom, L Stolt. Sol.

Energy Mater. Sol. Cells 41–2:345, 1996.68. R Bayón, C Guillén, MA Martínez, MT Gutiérrez, J Herrero. J. Electrochem. Soc.

145:2775, 1998.


69. R Bayón, C Mafftiotte, J Herrero. Thin Solid Films 353:100, 1999.70. R Bayón, J Herrero. Appl. Surf. Sci. 158:49, 2000.71. R Bayón, J Herrero. Thin Solid Films 387:111, 2001.72. T Yoshida, K Yamaguchi, H Toyoda, K Akao, T Sugiura, H Minoura. Proc. Elec-

trochem. Soc. 97–20:37, 1997.73. P Pramanik, MA Akhter, PK Basu. Thin Solid Films 158:271, 1988.74. CD Lokhande, A Ennaoui, PS Patil, M Giersig, M Muller, K Diesner, H Tributsch.

Thin Solid Films 330:70, 1998.75. CD Lokhande, KM Gadave. Turkish J. Phys. 18:83, 1994.76. P Pramanik, RN Bhattacharya. J. Mater. Sci. Lett. 8:781, 1989.77. P Pramanik, S Bhattacharya. Mater. Res. Bull. 25:15, 1990.78. KC Mandal, O Savadogo. J. Mater. Chem. 1:301, 1991.79. P Pramanik, S Biswas. J. Solid State Chem. 65:145, 1986.80. SS Dhumure, CD Lokhande. Thin Solid Films 240:1, 1994.81. AJ Varkey. Sol. Energy Mater. 21:291, 1991.82. SS Dhumure, CD Lokhande. Mater. Chem. Phys. 27:321, 1991.83. SS Dhumure, CD Lokhande. Sol. Energy Mater. Sol. Cells 28:159, 1992.84. GA Kitaev, TP Bol’shchikovsa. Izv. Akad. SSSR, Neorg. Mater. 2:65, 1966.85. GA Kitaev, TP Bol’shchikovsa, TA Ust’yantseva. Izv. Akad. SSR, Neorg. Mater.

3:1080, 1967.86. TP Bol’shchikovsa, GA Kitaev, VI Dvoinin, MV Degtyarev, LM Dvoskina. Izv.

Akad. SSSR, Neorg. Mater. 16:387, 1980.87. H Meherzi-Maghraoui, P Cowache, D Lincot, M Dachraoui. J. Chim. Phys. 96:259,

1999.88. H Meherzi-Maghraoui, M Dachraoui, S Belgacem, KD Buhre, R Kunst, P Cowache,

D Lincot. Thin Solid Films 288:217, 1996.89. MJ Mangalam, KN. Rao, N Rangarajan, CV Suryanarayana. Brit. J. Appl. Phys.

(J. Phys. D), Ser. 22:1643, 1969.90. MJ Mangalam, KN. Rao, N Rangarajan, CV Suryanarayana. Ind. J. Pure and Appl.

Phys. 7:628, 1969.91. SS Dhumure, CD Lokhande. Mater. Chem. Phys. 28:141, 1991.92. M Ristova, P Toshev. Thin Solid Films 216:274, 1992.93. AA Velykanov, EK Ostrovskaya, NP Garina, VA Turacova, AA Tchurkan. Ukr.

Chim. Zh. 49:764, 1983.94. B Pejova, M Najdoski, I Grozdanov, SK Dey. Mater. Lett. 43:269, 2000.95. AB Kulkarni, MD Uplane, CD Lokhande. Thin Solid Films 260:14, 1995.96. A Mondal, P Pramanik. Thin Solid Films 110:65, 1983.97. MJ Mangalam, KN Rao, N Rangarajan, CV Suryanarayana. Jpn. J. Appl. Phys.

8:1258, 1969.98. C Puscher. Dingl. J. 195:375, 1870.99. CD Lokhande. J. Phys. D: Appl. Phys. 23:1703, 1990.

100. RD Engelken, HE McCloud, C Lee, M Slayton, H Ghoreishi. J. Electrochem. Soc.134:2696, 1987.

101. P Pramanik, PK Basu, S Biswas. Thin Solid Films 150:269, 1987.102. MTS Nair, PK Nair. Semicond. Sci. Tech. 6:132, 1991.103. AJ Varkey. Int. J. Mater. Prod. Technol. 12:490, 1997.


104. PL Allen, A Hickling. Chem. Ind. 51:1558, 1954.105. T Jiang, GA Ozin, RL Bedard. Adv. Mater. 6:860, 1994.106. PK Nair, MTS Nair. J. Phys. D:Appl. Phys. 24:83, 1991.107. MTS Nair, PK Nair. J. Phys. D:Appl. Phys. 24:450, 1991.108. D Hariskos, R Heberholz, M Ruckh, U Ruhle, R Schäffler, HW Schock. In: 13th

ECPV Solar Energy Conf., Nice, France, 1995, p 1995.109. M Ristov, G Sinadinovski, I Grozdanov, M Mitreski. Thin Solid Films 173:53,

1989.110. P Pramanik, RN Bhattacharya. J. Mater. Sci. Lett. 7:1305, 1988.


7Oxides and OtherSemiconductors

Most of the compounds deposited by CD have been sulphides and selenides. Apartfrom a very few examples of tellurides (and some related telluride experiments)and with a very few exceptions, discussed at the end of this chapter, what is left isconfined to oxides (including hydrated oxides and hydroxides and two examplesof basic carbonates.) This chapter deals mainly with these oxides. In addition, asnoted in Chapter 3, there are a number of slow precipitations that result in precip-itates, rather than films, of various other compounds, not necessarily semicon-ductors in the conventional sense. These potential CD reactions, briefly discussedin Chapter 3, will be somewhat expanded on in this chapter.

Oxide films are often deposited because of their electrical (resistance) and op-tical properties. A selection of such properties of CD oxides is given in Table 2.2.

The reader is strongly urged to read Section 3.2.4 (precursors for oxide de-position) before reading this chapter or at least to refer back to it when necessary.

7.1 GENERAL CONSIDERATIONS FOROXIDE/HYDRATED OXIDE/HYDROXIDEDEPOSITION

The old analytical chemistry literature is rich with methods involving homoge-neous precipitation from solution, the purpose being to obtain dense (thereforeeasily filterable), contamination-free precipitates for purposes of analyses. The


urea method, in particular, has been extensively used in the past to form precipi-tates of oxides and basic salts for analytical purposes. Urea hydrolyzes to ammo-nia and (bi)carbonate, and the ammonia hydrolyzes further to give OH� ions, witha subsequent increase in pH. This leads to the formation of hydroxides, hydratedoxides, carbonates, and basic salts.

The formation of films has often been noted in precipitations using urea.Thus, Gordon [1] noted: “The precipitation of basic salts with urea is character-ized by the formation of thin transparent films of precipitate which strongly ad-here to glass surfaces.” Also, in Ref. 2 Gordon et al. wrote (p. 39): “Basic thoriumformate adheres tenaciously to glass surfaces in the manner characteristic of thebasic salts precipitated by the urea method.” Basic sulphates of Al [3] and Ga [4],which under suitable conditions contain very little sulphate (and are probably ox-ides or hydroxides), have been observed to form on glass using urea precipitation.

This film formation was an undesirable side effect—for accurate analysis,the film needed to be removed and added to the precipitate. However, as pointedout by Gordon [1], “To those who have worked with the urea method, the exis-tence of these films will always be a reminder of Willard’s* fond hope that he willsomeday find a way of making all the precipitate adhere to the beaker so that itwill only be necessary to dry and weigh the beaker after discarding the solution.”

Film formation is, in retrospect, not surprising, since the slow reaction char-acteristic of many homogeneous precipitation reactions is normally required(among other factors) for appreciable film formation to occur. Furthermore, thefilms tended to be very adherent. In their book Precipitation from HomogeneousSolutions [2], Gordon et al. write on the films formed by precipitation of basic tinsulphate: “The films cannot be removed by scraping with a policeman. However,by adding a few milliliters of hydrochloric acid . . . the films are easily dissolved.”(For those readers who, like the author, found this first sentence evoking amusingmental images, a policeman is (or was) a glass rod with a piece of rubber attachedto the end, used to scrape deposits out of reaction vessels.) A notable exception tothis “easy” dissolution was deposition of “basic stannic sulfate” using urea. Thismaterial formed such an adherent film that “it poses a difficult removal problem”[2,4a]. The films could be peeled off in relatively large, transparent sheets bywarming for 30 min in a solution of (NH4)2SO4 and NaOH at pH of 9 � 0.5, a factthat suggests using CD for the preparation of thin, self-supporting films. (This alsomeant that the film particles adhered strongly not only to the substrate, but also toeach other.) Actually, though not analytically well defined, this “basic sulfate” ac-tually contained very little sulphate and was probably mostly SnO2 (see the be-ginning Sec. 7.2.14 on SnO2). It should be remembered that most of these prod-ucts form white precipitates; therefore, if nonscattering, the films may not even be

* H. H. Willard was one of the pioneers of homogeneous precipitation for chemical analysis.


visible. Thus film formation may have occurred in cases where it was not reportedor even recognized.

As a general point, it might be expected that the product of reaction betweenmetal ions and hydroxide is a hydroxide (or basic salt) rather than an oxide. Inmany reported cases, oxides are formed directly. This is probably due to two fac-tors. Many of the metal ions used (e.g., Pb, Sn, Tl, Ti, Zr, Si) do not readily, if at all, form simple hydroxides; most of these cations have a greater tendency toform what may be called oxide polymers, involving condensation to chains ofOMMMOMMMO species at the reaction pH. Some hydroxides (e.g., Ag, Cu,Mn) are not very stable and are quite readily converted to the oxide, even in aque-ous solutions. In some cases, simple hydroxides do form and need to be heated todehydrate to the oxide.

Very acidic (high valent) cations will readily hydrolyse in aqueous solution,often even at low pH. These cations tend to form the polymeric metal oxide chainsmentioned previously. This hydrolysis can be controlled by addition of boric acid(see Sec. 3.2.4.4) and forms the basis of a technique referred to as liquid phase de-position. This method can be reasonably included in the more general term ofchemical solution deposition, and is treated, although not comprehensively, inthis book. Ref. 5 deals more thoroughly with this technique and describes manycases of SiO2 as well as some examples of several other oxides not covered in thischapter.

7.2 SPECIFIC OXIDES AND HYDROXIDES

7.2.1 Antimony Oxide (Sb2O3)

Films of Sb2O3 (more strictly, Sb4O6) have been deposited from a room-tempera-ture solution of potassium antimonyl tartrate and sodium selenosulphate [6]. Thefilms showed a clear XRD pattern, and compositional analysis confirmed thecomposition. It is interesting that the selenide was not formed from this solution(it was formed if the Sb solution was mixed with triethanolamine and ammoniabefore adding the selenosulphate; see antimony selenide in Chap. 6). The mostlikely explanation for this is that the more alkaline solution, containing tri-ethanolamine and ammonia, keeps the oxide, which might tend to form, in solu-tion, both because Sb4O6 is soluble in sufficiently strong alkali and because ofcomplexation by the triethanolamine and ammonia. Hydroxide would be presentin much larger concentrations than selenide, even under mildly alkaline condi-tions. The resistivity of the films was of the order of 109 �-cm.

7.2.2 Cadmium (Hydr)oxide (Cd(OH)2, CdO)

Films of what was presumably Cd(OH)2 were deposited by heating to 80°C an al-kaline cyanide solution of Cd2� containing H2O2 [7]. After heating the white as-


deposited films at 250–300°C, they turned brown to give CdO with resistivities oftypically a few k�/sq. (CdO is normally a degenerate semiconductor with a lowresistivity.)

H2O2 was used in two other studies to deposit cadmium hydroxides or hy-drated oxides. A mixed CdMOMOH film was deposited from a Cd-ammine so-lution at pH � 10 and at various temperatures onto glass and quartz [8]. Thedeposited films exhibited clear XRD peaks that were identified withCd(O2)0.88(OH)0.24 and were (111) textured. However, if KBr was added to the de-position solution, the preferred texture became (200) and the rate of film growth,as well as the terminal thickness, decreased. Annealing the films at ca. 200°C con-verted them into CdO with the same texture present as in the as-deposited films.From the optical spectra, a bandgap of ca. 2.6 eV was obtained (slightly higher forthe Br-free bath and slightly lower for the Br-containing bath). The resistivity ofthe films varied from 2 to 20 k�/sq (2 � 10�2 �-cm). In this case, thicker filmshad a higher resistivity, and this was ascribed to cracking of the thicker films, ob-served in optical microscopy studies. The second study used a similar solution andobtained films, identified by XRD as CdO2, with a high resistivity (106–107 �-cm) that converted to low-resistivity (10�3 �-cm) CdO on annealing in air at450°C with a direct bandgap of 2.3 eV [9].

Deposition without H2O2 has also been described, using ammonia-com-plexed Cd. In one study [10,11], deposition was carried out at either room temper-ature or 50°C. The as-deposited Cd(OH)2 films on glass were annealed at 400°C ineither air or an inert atmosphere to convert them to CdO. The as-deposited film wasX-ray amorphous, while the annealed film was polycrystalline CdO. A bandgap of2.2 eV (a little lower than the standard value of 2.4 eV) was obtained from the op-tical spectra of the CdO films, and the transmission in the nonabsorbing region washigh (up to 90%). The resistivity of the as-deposited Cd(OH)2, as expected, washigh (107 �-cm), while that of the CdO was 10�3 �-cm. In the other study, a higherammonia concentration was used and deposition was carried out at room temper-ature [12]. It was noted that the uniformity of the films was better than when de-posited at higher temperature. The Cd(OH)2 was heated in air at 150°C to convertit to CdO (the previous films required heating to ca. 275°C to convert the Cd(OH)2

to CdO [11]). (This is a not insignificant difference for what is assumed to be iden-tical material.) The bandgap was 2.3 eV, and, as in the previous study, the filmswere very transparent to photons with less energy than the bandgap. The resistiv-ity of the films was between 2 � 10�2 and 5 � 10�2 �-cm.

7.2.3 Cobalt (Hydr)oxide and Hydroxy-Oxide(Co(O)OH, Co(OH)2, CoO, Co3O4)

Co(O)OH was deposited from an ammonia-complexed solution of CoCl2 [13].The Co(II) ammonia complex was allowed to oxidize for two days in air to the


more stable trivalent cobaltic (III)–ammine complex. Heating the cobaltic–am-mine complex to 65–70°C resulted in deposition of ca. 0.1 �m of an adherent,brown CoO(OH) deposit on a glass substrate after 4 hr. Spectral measurements al-lowed an estimation for the bandgap of this film of ca. 2.4 eV. The Co(O)OH wasoxidized to an adherent Co3O4 film by heating over 300°C in air.

Films of CoO were deposited from a somewhat similar bath (but not left tooxidize) after annealing in O2 at 350°C; Co3O4 started to appear only at 500°C[13a].

Hydrolysis of urea to increase the pH by formation of ammonia was used todeposit Co(OH)2 by heating a solution of Co2� with urea at 100°C [14]. The pinkfilm was shown by XRD to be a mixture of phases of Co(OH)2. Heating the filmsat 350–400°C converted the hydroxide to Co3O4.

7.2.4 Copper Oxide (Cu2O)

Cu2O films were deposited by treating a thiosulphate-complexed solution ofCu(NO3)2 with NaOH [15]. This was based on an early study where a glass sub-strate was alternately and repeatedly dipped in NaOH and then Cu-thiosulphatesolutions [16]. The thiosulphate (S2O3

2�) both reduced the Cu(II) to Cu(I) andacted as a complexing agent. The films were deposited at 60–70°C, resulting in athickness of ca. 0.3 �m in one hour. The substrates—glass slides or polyesterfilm—were precoated with very thin CuxS films by immersion in copper thiosul-phate solution at 40°C. While the role of this prelayer was not clear, it was impliedthat in its absence, the Cu2O films were not uniform.

While the mechanism for the deposition was not discussed, the instability ofthe copper hydroxides (the hydroxide of Cu(I) probably does not even exist) to-ward dehydration, together with the reducing action of the thiosulphate, leads tothe expectation that Cu2O will be the product of the hydrolysis of Cu(I) in alka-line solution. It should be noted, however, that the Cu-thiosulphate solution itselfis not very stable and apparently forms predominantly CuxS in the absence ofNaOH.

X-ray diffraction showed the film to be Cu2O, with no detectable amount ofCuO and with a crystal size, estimated from the peak widths, of �20 nm. Opticaltransmission measurements of the films gave a value of (indirect) bandgap of 2.28eV (literature room-temperature bandgap �2.1 eV but is rather variable).

The electrical resistivity of the films (between 0.1 and 0.25 �m thick), mea-sured through Au contacts, was ca. 2.5 k�/sq (ca. 5 � 10�2 �-cm). This value in-creased with air-annealing (250°C, 20 min) up to 60 k�/sq. The relatively low re-sistivity was attributed to incorporation of S, either from the CuxS prelayer or byhydrolytic decomposition of S2O3

2� to S2�. Treatment of the films with Na2S so-lution decreased the resistivity by nearly two orders of magnitude, and S wasfound in the films. It is likely that the surface of the Cu2O crystals was partially


converted to CuxS, and surface conduction via this surface layer was responsiblefor the enhanced conductivity.

7.2.5 Indium Oxide (In2O3)

In2O3 was deposited from a solution of InCl3 that was slowly hydrolyzed to formthe hydroxide [17]. The rate of hydrolysis was slowed down sufficiently to pre-vent rapid bulk precipitation by a combination of a freezing agent (sodium citrate)and a relatively low pH (7.5). Ag� (as AgNO3) was added, supposedly as a cata-lyst (although it is not clear what needed to be catalyzed) and to improve adher-ence. The citrate will also act as a complexant for the In3�, which may be an im-portant factor. Addition of SnCl4 to the deposition solution allowed doping withSn [tin-doped indium oxide or indium tin oxide (ITO)]. The film grew to a termi-nal thickness of ca. 360 nm in 30 min. At higher solution pH, the terminal thick-ness decreased. Heating at 200°C for 2 hr in a vacuum resulted in conversion ofthe hydroxide(s) to crystalline oxide, with an average grain size of ca. 25 nm(In2O3) and ca. 54 nm (ITO).

Figure 7.1 shows the optical transmission and reflectance spectra of the twofilms. The main difference is an increase in the mid-IR reflectance of the dopedfilm compared with the undoped one, due to the high free electron concentration

FIG. 7.1 Transmittance and reflectance spectra of In2O3 films. Broken lines: undopedIn2O3; solid lines: ITO (10% Sn). (Adapted from Ref. 17.)


in the conduction band. The bandgaps measured from the spectra were ca. 3.5 eV(undoped) and ca. 3.6 eV (ITO).

Electrical measurements of the In2O3 (ITO) films gave temperature-inde-pendent resistivities (�) of 2 � 10�2 (10�3) �-cm, carrier concentrations (Nd) of2 � 1020 (1021) cm�3, and mobilities (�) of 3 (17) cm2V�1sec�1. The tempera-ture independence of the resistivity indicated that the films, even the nominallyundoped ones, were degenerate semiconductors.

In2O3 has been deposited on Sn/Pd-activated glass by first depositing a filmof In(OH)3 and then heating in air at a temperature of 200°C or more [18]. TheIn(OH)3 was deposited using a solution of dimethylamineborane and indium ni-trate maintained at 60°C. The deposition rate was dependent on the borane con-centration up to a limiting concentration of 0.03 M, and the film thickness wasproportional to deposition time, with final thicknesses of ca. 1 �m. X-ray diffrac-tion showed mainly one sharp peak corresponding to the (210) plane of In(OH)3,which, after annealing, converted to polycrystalline In2O3.

Optical transmission spectroscopy of the In2O3 film showed a high trans-mission at 800 nm, gradually decreasing with decreasing wavelength, character-istic of a somewhat scattering film. The bandgap was estimated from the spectrumto be 3.6 eV.

Electrical conductivity measurements of the as-deposited In(OH)3 showedan expectedly high resistivity of ca. 109 �-cm. That of the annealed oxide film de-creased to 33 �-cm (carrier concentration � 1.85 � 1016 cm�3; mobility � 10cm2V�1sec�1). The resistivity is high compared to many other In2O3 films (whichare often used as transparent conductors), mainly due to the low carrier concen-tration, implying a high degree of stoichiometry.

Finally, although no attempt was made to convert the film to oxide,In(OH)3, for use as a buffer layer on PV cells (see Chap. 9), was deposited from athiourea-based solution of InCl3 at a pH of 3.3 [19]. Apparently no sulphide wasformed, possibly due to the relatively high (for In) pH, which favored hydroxideformation.

7.2.6 Iron Oxides and Hydroxy-Oxide

The instability of Fe(III) compounds toward hydrolysis has been exploited to formFe(O)OH films [20]. The substrate in this study was a sulphonated-vinyl termi-nated self-assembled monolayer (SAM). Deposition was accomplished by heatingFe(NO3)3 solutions. The pH of the solution was rather critical; a pH of 2.0 orslightly higher was necessary. At lower values of pH, hydrolysis did not occur; atappreciably higher values, rapid hydrolysis occurred, resulting in precipitationrather than film deposition [e.g., at a pH of ca. 3, only very thin films (ca. 5 nmthick) of colloidal Fe(O)OH particles formed]. The films were columnar, with acolumn diameter of ca. 20 nm, and the columns were composed of lamellae ca. 2


nm in size. Nucleation of the (noncolloidal) film occurred by binding of Fe speciesto the sulphonate endgroups (one Fe to two sulphonate groups, measured by XPS)[21,22]

�-Fe2O3 was deposited on Si (111) or Si (100) by using hydrolysis of ureaat high temperatures [23]. An aqueous solution of Fe(NO3)3 and urea (pH between5 and 6), together with the Si substrate was heated to between 100 and 200°C (pre-sumably in an autoclave) for 4–24 hr. X-ray diffraction showed the formation of�-Fe2O3 with some (101) texture. From SEM measurements, the films were ca.100 nm in thickness, with a morphology depending on the crystal face of the Si.For (111) Si, the grains were spherical and 10 nm in size; for (100) Si, columnargrains, 30 � 5 nm, were obtained.

If reduced Fe powder was added to the preceding solution at an optimum pHof 6–7, magnetite (Fe3O4) was deposited onto Si (100) or �-Al2O3 at ca. 140°Cover several hours [24]. No other phase was found in the XRD spectrum. It wassuggested that the Fe3O4 formed by reaction between Fe(OH)3 (presumablyformed by hydrolysis of the ferric nitrate) and Fe(OH)2 formed by hydrothermaloxidation of the Fe powder. Particle sizes of 150 nm (on Al2O3) and 50 nm (on Si)were measured by SEM.

By adding Co(Cl)2 to the deposition solution and heating the resulting filmsin air (2 hr at 400°C) conversion of the magnetite to �-Fe2O3, doped with Co, oc-curred [25]. The grains were needle-shaped (50 � 10 nm). The films exhibitedgood magnetic properties.

A variation of the foregoing urea method was used to deposit �-Fe2O3 onSnCl2-sensitized glass substrates [26]. A solution of FeCl2 and urea at pH � 3 washeated at 90°C for 2 hr. The as-deposited film was probably FeO(OH) (the hy-droxide group was seen in FTIR studies). On annealing at 350–400°C (presum-ably in air), �-Fe2O3 was formed, with a crystal size of 22 nm. Optical spec-troscopy of the as-deposited film showed a direct bandgap of 3.2 eV and a weak(possibly indirect) absorption starting at ca. 2.2 eV. The bandgap of the oxidizedfilms was 2.0 eV. Resistivity was ca. 2 �-cm, which dropped by a factor of up tofour when exposed to high humidity, suggesting possible use as a sensor for wa-ter vapor.

Magnetite was also deposited on glass by the dimethylamine borane tech-nique described for In2O3, using a solution of Fe(NO3)3 and dimethylamine bo-rane with a pH � 3.5 at 20°C [27]. At higher deposition temperatures, Fe(O)OHand Fe2O3 were apparently also formed, and the magnetite content decreased, un-til, at 60°C, no magnetite was observed and the films were yellow (the magnetitefilms were black). The formation of pure Fe(III) oxide (hydroxide) at higher tem-peratures, compared to the mixed Fe(II)/Fe(III) magnetite at lower temperatures,is likely due to more facile oxidation of Fe(II) to Fe(III) at higher temperature. Theborane is a reducing agent, and this is presumably the reason that the mixed va-


lency magnetite, which can be envisaged as Fe(III)2O3�Fe(II)O, is formed at lowertemperature. The resistivity of the magnetite was ca. 2 k�-cm—much higher thanusual for magnetite. This was explained by a lower concentration of Fe(III) thanexpected from stoichiometry, which was obtained in these films. Magnetic prop-erties of the films were described.

7.2.7 Lead Oxide (PbO2)

A solution of Pb2� ions can be oxidized to PbO2 by persulphate [28]:

Pb2� � S2O82� � 2H2O → PbO2 � 2SO4

2� � 4H� (7.1)

This reaction normally resulted in a precipitate of PbO2. However, this processwas subsequently modified to give films of PbO2 without precipitation in solu-tion, apparently based on the observation of the authors that films of PbO2 weresometimes observed to form on the walls of the glass beaker using this reaction[29].

The film deposition was carried out at room temperature from an aqueoussolution of plumbous acetate, ammonium acetate, and ammonium persulphate, us-ing NH4OH to bring the pH to 6. A trace of AgNO3 was added as a catalyst for re-action 7.4 [29]. A film of PbO2 ca. 50 nm thick was formed in an hour. Once thisinitial film was deposited, thicker films could be built up, usually at a somewhathigher pH, in the absence of the AgNO3. The initial film formation appears to bea pure CD reaction. However, electrochemical studies of further film buildupshowed that an electroless deposition mechanism, involving two partial electro-chemical reactions, was responsible for film formation.

A few words on the difference between CD and electroless deposition are inorder here. Electroless deposition is related to electrodeposition, except that in-stead of charge being supplied by an external power supply, it is supplied inter-nally by oxidation (reduction) of a strong reducing (oxidizing) agent; the two par-tial electrochemical reactions occur at different sites on the substrate (the substrateis initially a sensitized solid and subsequently the deposited material itself). Thisimplies a reasonable electrical conductivity of the material to be deposited. Forthis reason, electroless deposition is used mainly for metals, but can be used forelectrically conducting compounds, of which PbO2 is an example. Since chargetransfer is involved, a change in valence state of the metal cation normally occursbetween solution and film.

Mindt [29] described some properties of these films (thicker electrolessfilms, not the initial purely CD ones). Electron diffraction showed that the filmwas �-PbO2. The crystal (more correctly the particle) size was found, by electronmicroscopy, to be ca. 200 nm. The carrier density, measured by the Hall effect,was ca. 1021 cm�3. The resistivity was somewhat dependent on the pH of deposi-


tion, varying from 2 � 10�3 (pH 7) to 3 � 10�2 �-cm (pH 10), although the car-rier density was not found to vary appreciably with pH (implying that the mobil-ity did vary with pH).

In a similar study, the deposition conditions were modified [pH of 8 (by am-monia), no Ag� catalyst, and a deposition temperature of 80°C] [30]. A primarythin film was deposited, followed by a second deposition, resulting in films sev-eral microns thick. Optical absorption spectroscopy gave a bandgap of 1.7 eV. Thefilm resistivity was 1.3 � 10�3 �-cm (carrier density � 8 � 1019 cm�3; mobility� 50 cm2V�1sec�1).

White films of 6PbCO3�3Pb(OH)2�PbO (from XRD analysis) were slowlyformed over a few days from alkaline-complexed Pb2� solutions that contained acolloidal hydrated oxide phase and that were exposed to air [31]. This was due toreaction with CO2 in the air (see Sec. 5.3.3).

7.2.8 Manganese Oxide (Mn2O3, MnO2)

Aqueous solutions of permanganate will slowly oxidize, forming a brown film[not clear if this is Mn2O3, MnO2, or Mn(O)OH] on the walls of the vessel inwhich it is stored. Increase of either acidity or alkalinity of the solution can accel-erate this decomposition reaction. This reaction has been used to treat polymersubstrates prior to chemical deposition, for improved adhesion [32]. In that work,the MnMO film was dissolved before the CD process, and the improved adhesionwas probably due to some increase in the hydrophilic character of the mostly hy-drophobic polymers.

Mn2O3 films have been deposited on glass from an ammoniacal solution ofMn2� [33]. Ammonium chloride was added to decrease the pH and slow down therate of hydrolysis. The initial product was believed to be Mn(OH)2, which oxi-dized in air to Mn2O3.

The persulphate technique used for PbO2 described earlier was extended toMnO2 deposited on glass, using manganous acetate in place of lead acetate, to-gether with ammonium persulphate and AgNO3 as a catalyst [34]. Adherent filmsup to 0.5 �m could be obtained. No XRD pattern was found for the films, imply-ing that the deposit was amorphous or made up of very tiny nanocrystals (nonde-fected crystal size of 2 nm or more can usually be detected by careful XRD). Theresistivity of the films was very dependent on the solution pH, with values of 4 �102 �-cm (pH 8) and 2 � 104 �-cm (pH 6.3).

7.2.9 Molybdenum Oxide

An early attempt to deposit Mo—S on various metal substrates using ammoniummolybdate and thiosulphate resulted in films that were found to be sulphur freeand believed to be an oxide, although this was not investigated further [35].


7.2.10 Nickel Oxide (NiO)

A brief communication on the deposition of NiO using persulphate was described[36]. An ammoniacal solution containing NiSO4 and potassium persulphate wasused to deposit black NiO on glass at room temperature. It was suggested, basedon the importance of the NH4OH (NaOH or KOH did not give NiO; other alkalinereducing species, such as amines, did) that a mixture of NiO and Ni2O3 formedand that the higher-valent Ni2O3 was reduced by the ammonia

The NiO was confirmed by XRD. Optical absorption spectroscopy was usedto estimate a direct bandgap of 1.75 eV (see later). The films were p-type (ther-moelectric probe), with a resistivity of 105 �-cm.

In another study, an ammoniacal solution of Ni2� was heated at 60–80°C todeposit a green-gray film of what was reported to be NiO [37]. This was assumedto form via the hydroxide, although no structural or compositional characteriza-tion of this deposit (or of the annealed film) was given. The deposition was car-ried out in a beaker, and no deposition occurred at room temperature; this suggeststhat deposition occurred by loss of ammonia. Heating this film in air at 280°Cformed Ni(O)OH. From the optical transmission and reflection spectra of the NiOand Ni(O)OH, it appears that their absorption spectra were very similar, with aweak absorption (possibly also scattering) in the visible and strong absorption inthe near-UV region. The resistivity of the Ni(O)OH was 800 �-cm, while that ofthe as-deposited NiO was apparently too high to measure.

Ni(OH)2 was deposited from a solution of urea and Ni2� ions at an initialpH of 6 and a temperature of 100°C. The hydroxide was then annealed in air at350–400°C to convert it to NiO [38]. Films almost 1 �m thick were obtained af-ter ca. 2 hr, with an average crystal size (from XRD) of 13 nm. Optical absorptionspectroscopy of the annealed films gave a direct bandgap of 3.6 eV, somewhatlower than the rather variable literature values of 3.7–4.0 eV. It should be men-tioned that it is not as simple to correlate the band structure of NiO with its opti-cal and electrical properties as it is for most of the other semiconductors dealt withhere. This feature, common to many transition metal compounds, is a conse-quence of electrons in (often-narrow) d-bands that are relatively localized by elec-tron–electron repulsions. Thus, although the 3d band in NiO (which in the purestate is green, like the hydrated Ni2� ion) is only partially filled, pure NiO is in-sulating due to the localized 3d electrons. The conductivity (and black color) ofNiO as it is normally obtained is due to nonstoichiometry leading to doping. Theroom-temperature resistivity of the annealed NiO in this case was several M�/sq.(several hundred �-cm).

7.2.11 Silicon Oxide (SiO2)

SiO2 films were deposited on soda lime glass from a silica gel–saturated solutionof hydrofluorosilicic acid (H2SiF6) and boric acid [39]. The boric acid reacts with


the H2SiF6 to form SiO2 by removal of HF formed in the H2SiF6 hydrolysis equi-librium:

H2SiF6 � 2H2OD 6HF � SiO2 (7.2)

H3BO3 � 4HFD BF4� � H3O� � 2H2O (7.3)

The SiO2 formed gradually deposited on the substrate. The deposition rate wastypically 10–20 nm/hr, depending on the boric acid concentration and solutiontemperature. Film thickness was ca. 100 nm.

Infrared measurements showed that the films had a higher concentration ofSiMOMSi bonds than some other silica films made by different techniques. Thiswas interpreted to mean that the silica network of the films was more orderly, aproperty that was evidenced by greater stability of the films against chemical etch-ning and good blocking properties to sodium diffusion from the soda glass, com-pared to many other silica films.

This is the first of many SiO2 depositions using essentially the same tech-nique. The others are given in Ref. 5.

7.2.12 Silver Oxide (AgO and Ag2O)

Oxides of Ag have been deposited by deposition from a triethanolamine-com-plexed Ag solution at a pH � 11.5 [40]. At room temperature, a black depositformed over some hours that converted into a brown film when air-annealed at150°C. Based presumably on the color of these films, the as-deposited film wasassumed to be AgO, which turned into Ag2O on annealing (no structural charac-terization of the films was reported). From the optical spectra, a bandgap of 2.25eV was estimated for the annealed films ([literature (direct) bandgap of Ag2O �1.2 eV; a pure semiconductor with a direct bandgap of 2.25 eV should be orange].Both films were insulating (no values given) but became much more conductingwhen high voltages (�1 kV) were applied to two laterally spaced Ag electrodeson the (ca. 500-nm-thick) films.

Using a similar solution, films of either Ag or AgO were deposited on bothglass and polyester film [41]. Addition of triethanolamine to a Ag� solutioncaused initial precipitation (silver oxide or hydroxide), which redissolved in ex-cess triethanolamine. Deposition from a solution where some precipitate remainedresulted in AgO (possibly with some Ag2O), while a solution where this precipi-tate was completely redissolved gave metallic Ag. The reducing action of the freetriethanolamine present in the latter case may be the cause of the formation ofmetallic Ag.

7.2.13 Thallium Oxide (Tl2O3)

The same persulphate technique described earlier for PbO2 and MnO2 was alsoused to deposit Tl2O3, with thallous acetate in place of the other metal acetates


[34]. Adherent films could be deposited up to a thickness of 10 �m—very thickfilms for the CD method. X-ray diffraction confirmed that the films were cubicTl2O3. The film resistivity was 5 � 10�4 �-cm.

The modifications employed by Bhattacharya and Pramanik for the PbO2

deposition (higher pH and temperature, no Ag�) were also used for Tl2O3 [30].Film resistivity was similar to that of the previous study (3.7 � 10�4 �-cm),with a carrier concentration of 4 � 1020 cm�3 and a mobility of ca. 50cm2V�1sec�1.

7.2.14 Tin Oxide (SnO2)

Very strongly adhering films of “basic stannic sulfate,” which was probably SnO2

(the Sn:SO42� ratio was 44:1) were reported using an aqueous solution of SnCl4,

urea, H2SO4, (NH4)2SO4, and HCl at pH 0.5 [4a]. The films adhered to the wallsof the glass deposition vessel so strongly that HF was one of the few reagentsavailable to remove them. Films which are mainly (hydr)oxides of Al, Ga and Thhave also been reported in the early literature using the urea method [2]. Tin salts,in particular those of Sn(IV), are readily hydrolyzed, and the stable product isSnO2. This has been exploited in a number of studies, most with minor differencesbetween them, to deposit SnO2 films.

SnO2 films were deposited using SnCl4 and NH4F (the latter apparently asa complexing agent to slow down hydrolysis of the Sn4� by the alkaline solution)[33]. The resistivity of the as-deposited film was 200 �-cm (10�2 �-cm for Indoped).

A similar method was described a few years later in which the NH4F wasused as a freezing agent to slow down the rate of hydrolysis [42]. AgNO3 wasadded as a catalyst(?) and to improve the film adherence, although it is not clearwhy a catalyst was needed or even desirable, since the objective was to slow downthe reaction. NaOH was used instead of ammonia to adjust solution pH to between7.5 and 8.5.

The growth rate was linear and decreased with decreasing pH, with a limit-ing thickness that increased with decreasing pH. The growth rates and terminalthicknesses were similar to those for ZnO deposited by the same technique (seeFig. 7.2 in Sec. 7.2.17), only the rates were two to four times slower. These rela-tionships were explained in the same way as for ZnO.

The films were found to be SnO2 with the rutile structure (by XRD), with agrain size of 20–30 nm (by TEM). Optical transmission and reflectance spec-troscopy showed that the films were close to 80% transmitting up to ca. 1 �m andhighly reflecting in the mid-IR. These spectra were similar to those of ZnO shownin Figure 7.3, (Sec. 7.2.17), except for the absorption in the UV of the undopedZnO due to the lower bandgap. The bandgap (direct) was 3.56 eV. Electrical mea-surements gave a resistivity (�) of 0.1 �-cm, carrier concentration (Nd) of 1019


cm�3, and mobility (�) of ca. 6 cm2V�1sec�1. Annealing in vacuum decreased �to a minimum of 2 � 10�3 �-cm at 375°C.

Antimony-doped SnO2 films were deposited by adding SbCl3 to the depo-sition solution. Sb is a well-known n-type dopant used to increase the conductiv-ity of SnO2 films. The Sb concentration in the films increased linearly with that inthe deposition solution and was somewhat less than the solution concentration(e.g., 6% Sb in solution gave ca. 4% in the film). The Sb doping increased boththe visible/near-IR transmission and mid-IR reflectance of the films, compared tothe undoped films. These spectra are similar to those for doped ZnO (Fig. 7.3), andthe effect of doping can be explained in the same way. The bandgap increased to4.1 eV, compared to 3.56 eV for the undoped film, explained through band fillingby free electrons.

The conductivity, Nd, and � all increased with increasing Sb concentrationup to 5% and then decreased again with increasing Sb (Nd leveled off). This wasdue to the obvious increase in Nd with increased doping, the measured increase ingrain size with doping (30–65 nm), resulting in increased � followed by segrega-tion of dopant at the grain boundaries at greater Sb concentrations, which againdecreased �. The resistivity of the 5% Sb films decreased from 10�3 �-cm as de-posited to almost as low as 10�4 �-cm on annealing at 375°C.

In another preparation [43], ammonia gas was passed through a solutionof SnCl4, the precipitate rinsed well (to remove Cl, which caused the final filmsto be porous—a useful observation since porous films are sometimes preferredover compact ones), and redispersed in concentrated HNO3 to give a semitrans-parent sol at a pH of 5–7. Films were deposited from this solution onto Si (100)by heating at 60–100°C for 4 hr, followed by 100–200°C for 6–12 hr. The ini-tial lower-temperature step was necessary to obtain nucleation; if omitted, nofilm was deposited. Films ca. 200 nm thick were obtained with a crystal size of 3.5–4 nm.

As with the original deposition of “basic stannic sulphate,” urea was used inanother deposition to slowly increase the pH of the initially strongly acidic (byHCl) solution, thereby hydrolyzing the SnCl4 [44]. Besides slowing down the hy-drolysis of the SnCl2 due to increased acidity, the HCl also complexed the tin ashexachlorostannate, (SnCl6)2�, further slowing down the hydrolysis. From the op-tical spectrum, a bandgap of 4.0 � 0.1 eV was estimated. These films were usedas buffer layers on CuInSe2 solar cells (see Chap. 9).

SnO2 was deposited on hydrolyzed Si and on Si coated with sulphonate-terminated self-assembled monolayers from a solution of SnCl4 in dilute HCl at80° [45]. The films, up to 65 nm thick and consisting of a dense-packed aggregateof SnO2 nanocrystals (5–10 nm) together with some amorphous basic tin oxide,contained ca. 3 at.% Cl. They were adherent on all substrates, although the adher-ence and homogeneity on Si was less reproducible than on the monolayer-coatedSi. Films were also deposited using a continuous flow system. The films were sim-


ilar to those using a static solution, but the deposition rate was considerably faster(�20 nm hr�1 compared to an average of several nm hr�1 for the static solution),and thicker films (160 nm) could be obtained.

Films of SnO2 up to 0.5 �m thick were grown by hydrolysis of aqueousSnF2 (a change from the usual SnCl4) solutions, optimally at ca. 60°C on variousglasses and on Si [46]. The as-deposited films contained 6–16 mole% F, an n-typedopant in SnO2. Films annealed at 300°C in air contained almost no F and werecrystalline SnO2 (cassiterite). It was suggested that the as-deposited material pos-sessed a tin–oxygen polymer structure, with tin–fluorine bonds substituting forsome of the tin–oxygen bonds. Conductivities of the order of 10�2 �-cm weremeasured for films annealed at 500°C.

Deposition of SnO2 onto sulphonated polystyrene at 40°C has been brieflydescribed in a review by Bunker et al., although without experimental details [47].The point was made that the deposition of SnO2 involves hydrolysis and conden-sation reactions involving poorly characterized species as opposed to precipitationof an ionic salt. Also, the deposition is very dependent on pH—a change in pH ofone unit can change the solubility of the products by four orders of magnitude. Inspite of these factors, which imply difficulty in depositing films in the absence ofbulk precipitation in solution, they apparently succeeded in achieving preferentialnucleation of SnO2 on the sulphonated polystyrene. The films were dense andcomposed of cassiterite, with a grain size of 4 nm.

7.2.15 Titanium Oxide (TiO2)

Sulphonate-terminated self-assembling monolayers on Si were used as substratesfor TiO2 deposition [48]. The deposition solution consisted of TiCl4 in 6M HCl at80°C (considerably more dilute HCl solutions resulted in immediate bulk precip-itation, while much stronger solutions were stable against hydrolysis and thereforeno deposition occurred.)

The films, ca. 50 nm thick, comprised small (2–4 nm) nanocrystals ofanatase TiO2, possibly in an amorphous matrix, and were uniform, adherent, andpore free. In contrast, only a small amount of irregular deposit was formed on bareSi. The role of the sulphonate endgroups was believed to promote nucleation ofthe nanocrystals and/or facilitate attachment of TiO2 clusters in solution to thesubstrate. Hydrolysis of TiCl4 proceeds through various titanium hydroxy andchloro-hydroxy complex cations. The anionic sulphonate groups could thus pro-mote attachment and nucleation of these cationic complexes.

Annealing in air increased the crystal size (up to 25 nm at 600°C) withoutdamage to the film while retaining the anatase structure (the rutile structure wasbarely noticeable after 2 hr at 600°C).

Using this deposition technique, TiO2 was deposited onto patterned self-assembled monolayers [49]. Thioacetate-terminated trichlorosilane monolayers


were self-assembled onto oxidized Si substrates. Illumination through a grid-pat-terned mask with grid openings of ca. 10 �m resulted in photolysis of the ex-posed, somewhat hydrophobic thioacetate groups to hydrophilic sulphonategroups. TiO2 was then deposited from the TiCl4 solution onto the irradiated hy-drophilic regions of the substrate (see Figs. 2.5 and 2.6). The resistivity of thisTiO2 was ca. 109 �-cm.

TiO2 was deposited on (100) Si from a solution of TiO42� prepared by dis-

solving Ti metal in an aqueous solution of ammonia and H2O2 [50]. An initial low-temperature–final high-temperature regime similar to that described earlier forSnO2 by the same group was used at a pH between 6 and 7. As before, the low-temperature stage was necessary for film formation. A film formed in the low-temperature stage but was X-ray amorphous. After the high-temperature stage,pure anatase-phase TiO2 was obtained in the form of square platelets ca. 10 � 10nm in size. The film was highly (112) textured.

7.2.16 Vanadium Oxide

Vanadium oxide films were deposited by dissolving V2O5 in aqueous HF and im-mersing an Al plate in the solution (the Al acts as a scavenger for F� in place ofthe more commonly-used boric acid) [50a]. The brown film was X-ray amorphousbut crystallized on heating in air through a mixed V(IV)- V(V) oxide to V2O5 (inan inert atmosphere, VO2 was formed on annealing). The as-deposited materialwas believed to contain mainly V(IV).

7.2.17 Yttrium Oxide (Y2O3)

Basic yttrium carbonate [Y(OH)CO3] was deposited by CD and subsequently an-nealed in air at 600°C to Y2O3 [51]. Si wafers and self-assembled monolayers withsulphonate endgroups were used as substrates. An aqueous solution of YNO3 andurea was heated at 80°C in sealed vials. The increase in pH, together with gener-ation of carbonate from hydrolysis of urea (Sec. 3.2.4.1), resulted in formation ofthe basic carbonate.

The film thickness was 35 nm. Various analytical techniques were used toconfirm that the deposit was amorphous Y(OH)CO3. Annealing at 600°C was nec-essary to convert the film to crystalline Y2O3 (amorphous oxide was formed be-tween 300 and 400°C), with a film thickness of 25 nm and crystal size of ca. 20nm.

7.2.18 Zinc Oxide (ZnO)

ZnO is the most studied of all the oxides deposited by CD. This is largely due toits use as a transparent, electrically conducting layer.

The first description of CD ZnO arose from the observation that depositionof mixed (Cd,Zn)S resulted in large amounts of ZnO, and this led to development


of a technique for depositing pure ZnO [7]. The method is similar to that used forCdO—heating an alkaline cyanide solution of a zinc salt to 80–90°C (although, incontrast to the Cd case, no H2O2 was added). The films, a few hundred nanome-ters in thickness, were shown by XRD to be ZnO. Sheet resistivities were ca. 108

�/sq, which dropped three orders of magnitude after heating at 350°C for 10 minin forming gas.

ZnO was deposited by CD using the freezing technique described previ-ously for In2O3 [52]. The substrates (glass or quartz) were immersed in an aque-ous solution containing ZnCl2, NH4F (as freezing agent—probably also acting asa complexing agent for the Zn2�), and Ag� catalyst (as for In-O and SnO2, it isnot clear for what purpose), and NaOH was then added to a pH of between 7.5 and8.5. The resulting Zn(OH)2 films were then annealed for a few hours in air or vac-uum at 180–200°C. The rate of deposition and terminal thickness of the ZnO filmsare shown in Figure 7.2. The films grow faster at higher pH but with a lower ter-minal thickness than at lower pH. At a pH substantially lower than 7.5, no growthoccurs; the OH� concentration is too low to precipitate Zn(OH)2. As the pH is in-creased, the rate of formation of Zn(OH)2 increases. This results in increasing ho-mogeneous precipitation in the solution, leading to loss of reactant and thinnerfilms, until, at a pH substantially greater than 8.5, essentially all the Zn is precip-itated and no (or, more likely, only an ultrathin) film formation occurs.

Al-doped ZnO films were also deposited by adding AlCl3 to the depositionsolution. The amount of Al in the films (given as at.% with respect to the Zn con-centration) was somewhat smaller than that in the deposition solution but was pro-portional to the concentration in solution (up to the maximum measured concen-tration in the films of 5.5%).

The films (both ZnO and ZnO:Al) were wurtzite structure with a preferen-tial texturing (c-axis ⊥ substrate). No Al2O3 was found in the XRD spectra, sug-gesting either dispersal of the Al in the ZnO matrix or its presence as very tinycrystals of (hydr)oxide on the ZnO surface. TEM measurements showed an aver-age grain size of 25 nm (ZnO) and 45 nm (ZnO:Al).

Due to the possible application of these films for transparent electricallyconducting or infrared-reflecting purposes, the optical and electrical properties ofthe films were the subject of careful study. Figure 7.3 from 268 shows the opticaltransmittance and reflection spectra of both undoped and Al-doped films. Thedoped films have a higher visible/near-IR transmittance. The transmittance cutoffis blue-shifted for the doped film due to conduction band filling by electrons fromthe Al dopant; this results in an increase in the effective optical bandgap (theBurstein–Moss shift) and therefore an absorption blue shift. (The bandgaps mea-sured from the spectra were 3.40 and 3.98 eV for the undoped and doped films, re-spectively. The value of 3.4 eV for the undoped sample is considerably higherthan the usual value of 3.2 eV; these films were already quite highly conducting—see the electrical measurements later and Table 7.1). The mid-IR reflectance of thedoped film also increased due to free-electron reflection.


FIG. 7.2 ZnO film thickness vs. deposition time for different values of solution pH.(Adapted from Ref. 52.)

FIG. 7.3 Transmittance and reflectance spectra of ZnO films. Broken lines: ZnO; solidlines: ZnO:A1 (4 at.%). (Adapted from Ref. 52.)


The resistivity of the films decreased with increasing A1 content up to 4%and then increased again to the maximum measured A1 content of 5.5%. This isdue to a combination of increase of carrier concentration at 4% A1, which then lev-els off, and increase in mobility at 4% A1 followed by a sharp decrease at 5.5% A1.The increase in carrier concentration is clearly due to the doping. Increase in mo-bility was linked to the increased grain size of the doped films, while the decreasein mobility (and conductivity) at doping levels � 4% A1 was attributed to grainboundary segregation, resulting in higher intergrain barriers. The relevant electri-cal parameters are listed in Table 7.1. Further decrease in resistivity was obtainedupon annealing (see Table 7.1). In all cases where maximum conductivity was ob-tained (350°C in vacuum or 225°C in oxygen), the decreased resistivity was duemainly to increased carrier concentration, although a moderate increase in mobil-ity was also measured. For vacuum-annealing, these two effects were explained byan increase in oxygen vacancy concentration and desorption of oxygen from grainboundaries (therefore decrease in grain boundary barrier), respectively. The reasonfor the increase in carrier concentration and mobility after annealing in oxygen at225°C is not clear, although the decrease in these parameters on annealing at highertemperatures in oxygen follows naturally from the foregoing explanation.

Hexagonal ZnO films were deposited on glass or SnO2/glass using a solu-tion of zinc acetate complexed with ethylenediamine and with the pH adjusted tobetween 10.5 and 11.0 with NaOH [53,54]. The formation of the films occurredonly under conditions where Zn(OH)2 was calculated to be present in the deposi-tion solution. Good-quality, adherent films were obtained only on glass and withinnarrow pH and composition ranges. Films could be obtained outside these condi-tions, but they were then poorly adherent. In general, the adherent films formedunder conditions where the deposition was slow (relatively low pH, relativelyhigh complex:Zn ratio).

TABLE 7.1 Variation of Electrical Parameters of ZnO Films with Al Doping andAnnealing Conditions

200°C 225°C oxygen 350°C vacuum

at.% Al N � � N � � N � �

0 0.67 7.7 12 2.5 10.4 2.5 2.4 10.2 2.54 9.0 11.9 0.60 19 16.7 0.21 18 16.7 0.205.5 9.4 1.5 4

200�C is the basic anneal that converts the hydroxide into the oxide. The films were then reannealed.The two conditions given here are those that give the most highly conducting films. Higher tempera-tures in oxygen decrease the conductivity, while higher temperatures in vacuum have no further effect.N: free electron concentration (� 1020 cm�3); �: mobility (cm 2V�1s�1); �: resistivity � 10�3 �-cm.(From Ref. 52.)


It was proposed that deposition of adherent films proceeded on Zn clustersbound via OH groups to the glass surface. The poorer adherence on a SnO2 sur-face was explained by a lower concentration of hydroxy groups on this surface.

In a related study, McAleese and O’Brien showed how ZnO nucleated onglass and SnO2/glass from solutions of zinc acetate, ammonia, and thiourea (andsometimes also hydrazine) [55]. Such a deposition solution is normally used to de-posit ZnS, and this study showed that ZnO can form together with, or even insteadof, ZnS.

From the optical spectra of the films deposited from the ethylenediamine-complexed solutions, a bandgap of 3.15 eV was calculated (literature value 3.2 eV).

ZnO was deposited on Si (100) by heating an ammoniacal solution (at pH� 7) of ZnAc2 [56]. The heating regime was important: First the solution washeated to between 60 and 100°C for 6 hr and then the temperature was raised tobetween 100 and 200°C for 6–12 hr (in an autoclave). The initial lower tempera-ture was necessary to obtain deposition, and it was suggested that nucleation ofZnO on the substrate occurred during this step. No XRD pattern was obtained af-ter this lower-temperature stage; wurtzite ZnO was clearly seen by XRD after thehigh-temperature stage. The films, ca. 65 nm thick, were smooth, dense, and ho-mogeneous, with some (10.0) texture.

ZnO films for use as buffer layers in photovoltaic cells (see Chap. 9) havebeen chemically deposited from aqueous solutions of ZnSO4 and ammonia [57].The solution was heated to 65°C, and adherent, compact Zn(OH)2 � ZnO filmswere formed after one hour. Low-temperature annealing converted the hydroxideto oxide. The solution composition will be important in this deposition. On onehand, increased ammonia concentration will increase the pH and therefore the ho-mogeneous Zn(OH)2 precipitation in solution. However, further increase in am-monia concentration will redissolve the hydroxide as the ammine complex. Therewill clearly be an optimum ammonia (and zinc) concentration where Zn(OH)2

does form, but slowly enough to prevent massive homogeneous precipitation. Theuse of ammonia in (hydr)oxide deposition derives, in part at least, from its grad-ual loss by evaporation if the system is not closed [58]. Any open solution of anammonia-complexed metal ion (which forms an insoluble hydroxide or hydratedoxide) should eventually precipitate the (hydr)oxide for this reason alone.

The borane technique, described earlier for In2O3 preparation, was also usedto deposit ZnO on Sn/Pd-activated glass using a solution of Zn(NO3)2 anddimethylamine borane at 50°C [59]. The films were randomly oriented polycrys-talline ZnO, and the crystals were hexagonal shaped, with a typical size of 0.2 �m,at moderately high dimethylamine borane concentrations; at low concentrations,the grains were smaller and more irregular.

Optical transmission spectra gave an estimated bandgap of 3.3 eV. From thespectra, the films showed some scattering, with the most transparent films havingan approximate integrated transmission over the visible region of 70%, obtainedfrom a solution containing 0.05 M/l dimethylamine borane. This correlated with


the most regular morphology measured by SEM. A more recent study by the samegroup, using higher Zn concentrations and slightly higher deposition temperature(60°C), showed increased transmission up to nearly 80% over the visible rangewith a borane concentration of 0.1 M [60].

Electrical measurements of films as a function of boron doping were also re-ported in this study. The amount of boron in the films varied over only a smallrange, from 0.01% at very low borane concentrations in solution (a minimum con-centration was needed for the deposition to proceed) to 0.02% at a borane con-centration in solution of 0.1 M. The resistivity of the films was high in all cases,varying from ca. 20 k�-cm to 0.4 k�-cm with increase in boron content. Mobil-ity and carrier concentration measurements showed low values of both (maximumvalues: � � 1 cm2V�1 sec�1, N � 1.8 � 1016 cm�3). The low carrier concentra-tion implies that these films are highly stoichiometric. While unfavorable fortransparent conducting purposes (although annealing would probably improve theconductivity), the ability to make relatively insulating ZnO may be advantageousfor other purposes.

7.2.19 Zirconium Oxide (ZrO2)

A process similar to that described earlier for TiO2 was used to deposit ZrO2 films[61]. Self-assembled monolayers with terminal sulphonate or methyl groups on Siwere used as substrates; no film growth occurred on bare Si. The deposition solu-tion was Zr(SO4)2 dissolved in an aqueous HCl solution at 70°C.

The film growth was slow—15 nm after 4 hr and with a limiting thicknessof 40 nm after ca. 20 hr, although it was faster at the beginning of the deposition.The films were composed of a mixture of crystalline tetragonal ZrO2 and amor-phous material—probably a basic zirconium sulphate. The (thicker) films variedin their thickness, from up to 10-nm-sized crystals near the substrate to 2- to 3-nmcrystals, together with a greater proportion of amorphous basic sulphate, towardthe film surface. It was suggested that both electrostatic forces (between the neg-atively charged sulphonate surface groups and positively charged zirconium ox-ide and basic sulphate colloids) and van der Waals attractive forces cause the ob-served good adhesion between the films and sulphonate monolayers while theinferior adhesion to the uncharged methyl-terminated monolayers was due solelyto van der Waals forces.

Annealing the films for 2 hr at 500°C, though causing pyrolysis of themonolayer, did not damage the film (or its adhesion to the substrate), which re-mained tetragonal ZrO2 with a crystal size ca. 10 nm and with some sulphate, thelatter disappearing after prolonged annealing. Annealing at 600°C or higher re-sulted in a change to the monoclinic phase of ZrO2

Yttrium-doped ZrO2 was deposited by adding Y2(SO4)3 and urea and de-positing at 80°C (see deposition of Y2O3 described earlier). A higher pH (2.5–3.0)


was needed to cause coprecipitation of Y2O3, and the Y concentration in the filmwas ca. 25% of that in the deposition solution. Basic carbonates and sulphateswere also present in the films. Annealing in air at 500°C for 2 hr resulted in com-plete crystallization of the films to yttrium-stabilized zirconia.

7.2.20 Ternary Oxides

Deposition of two ternary oxides—(Cd,Sn)O and (Zn,Cd)O—will be mentionedvery briefly next. These will be treated more fully in Chapter 8. Also, a brief de-scription of a related technique (SILAR; see Sec. 2.11.1) that has been used forcomplex oxides and related compounds will be given.

7.2.20.1 Cadmium Stannate (Cd2SnO4)

The ammonium fluoride technique used by Raviendra and Sharma for ZnO (de-scribed earlier) has also been used by them to deposit cadmium stannate using amixture of CdCl2 and SnCl4 [52]. After annealing at over 200°C, Cd2SnO4 was ob-tained. Optical and electrical properties of these films are described in Chapter 8.

7.2.20.2 Zinc Cadmium Oxide (Zn,Cd)O

Films of ZnxCd1�xO, with varying values of x, were deposited from ammoniacalsolutions of Cd and Zn chlorides containing hydrogen peroxide at 45°C followedby annealing in air at 500°C, presumably to convert the hydroxides to oxide [62].The optical and electrical properties of these films are described in Chapter 8.

7.2.20.3 SILAR Deposition of Metal Oxides,Hydroxides, and Peroxides

In SILAR (successive ionic-layer and reaction) deposition, discussed in Section2.11.1, successive compound layers are built up from reaction between adsorbed(ideally mono-) layers and a reactive solution. This technique has been applied tosulphides, selenides, and oxides, including hydroxides, peroxides, and ternary ox-ides. As an example, films of LaxNbOy were deposited [63]. This technique wasused to deposit many other oxides, hydrated oxides, and peroxides (see referencesin Ref. 63). In view of the uncertain purpose of H2O2, often used in CD of oxides,the role of this chemical as explained in this work is of interest. Except at a pH be-low ca. 2, the surface of oxidized Si (Si was used as a substrate in these experi-ments) in aqueous solution, and also glass, is composed of SiMO� groups [see Eq.(2.16)]. These groups can attract metal cations to give an uncharged (for monova-lent cation) or positively charged (for a higher-valent cation) surface. Thus, for adivalent cation:

SiMO� � M2�D SiMOMM� (7.4)


This surface reacts with H2O2 to give

SiMOMM� � H2O2D SiMOMMMOOH � H� (7.5)

which, in neutral or alkaline solution, dissociates to give a negatively charged sur-face again:

SiMOMMMOOHD SiMOMMMOO� � H� (7.6)

Additionally, the ability of H2O2 to oxidize a metal to a higher-valent state, re-sulting in a more insoluble hydroxide (higher-valent metal hydroxides are moreinsoluble at a particular pH than the hydroxide of the metal in a lower-valent state)has been pointed out in this study.

7.3 OTHER SEMICONDUCTORS: Se AND SILVERHALIDES

The remaining semiconductors (apart from ternaries, which are treated in Chap. 8)that have been deposited by CD are elemental Se and silver halides. The little doneon these materials will be discussed here.

7.3.1 Se

Elemental Se exists in several forms, including amorphous (red), monoclinic(red), and gray (hexagonal). Gray Se is the most stable form. There are several re-ports on CD Se. The earliest is based on the fact that selenosulphate is unstable inacidic solution and, if made acid, will immediately precipitate red Se (note that itsanalogue, thiosulphate, behaves similarly but that a much lower pH is required toprecipitate elemental S). Se was deposited by slightly acidifying dilute (10–50mM) selenosulphate solution [64]. Films of amorphous, red Se were obtained at10–15°C, while gray Se was obtained at 30°C or higher. Using weak acids (citricor ascorbic), amorphous Se films were likewise deposited at 0°C [65]. These con-ditions slow the formation of Se enough to allow films (ca. 50 nm thick) to form.A direct bandgap of 2.0 eV was measured for these films. Heating at 85°C trans-formed the films to gray, hexagonal Se.

Another technique, which is not strictly true CD but close enough to war-rant inclusion, is photodeposition from an amorphous Se colloid [66–68]. The Secolloid was prepared by reduction of a solution of SeO2. Illumination of a sub-strate in this solution with light that was absorbed by the Se (bandgap 2.05 eV) re-sulted in film formation on a substrate. Film formation occurred in the absence ofillumination but was extremely slow, particularly at lower temperatures (at tem-peratures above 25°C, gray Se began to be formed). No XRD structure was foundfor the films (deposited below 25°C). Raman spectroscopy revealed the presenceof Se chains and rings in the film. The mechanism of the deposition was not com-pletely understood, but it was clearly connected with the photogeneration of


charges in the Se, since only superbandgap illumination was effective. Thebandgap of the films was measured, from the optical spectra, to be 2.05 eV [69].

7.3.2 Silver Halides (AgI, AgBr, AgCl)

Section 3.2.6 begins with the sentence “Although they do not appear to have beenused (at least deliberately) to form films, there are other slow anion-generating re-actions.” And near the end of this section is written “It should be stressed thatthese reactions were used to form precipitates and not films. There is no guaran-tee that films can be formed using these reactions. However, it is reasonable to ex-pect that, under the right conditions, it may be possible to produce films of thesecompounds. It is left as an exercise for the curious reader to find these ‘right’ con-ditions.”

For most of the period while this book was being written, the halides wereincluded in that section of Chapter 3. However, the very act of writing down thatit should be possible to produce films of compounds containing these anions ledme (after a couple of years, when the opportunity arose to spend an extended pe-riod of time in the laboratory during a visit to the university of Bern) actually totry to do this. For a number of reasons, halides were the anion of choice, specifi-cally silver halides. The slow generation of halides is discussed in Section 3.2.5and is based on the slow hydrolysis of haloalcohols.

AgI was deposited by hydrolysis of 2-iodoethanol (ca. 50 mM):

ICH2CH2OH � H2OD I� � H� � HOCH2CH2OH (7.7)

in an aqueous solution containing AgNO3 (ca. 10 mM) and a small amount of tri-ethanolamine (ca. 0.5 mM). The triethanolamine, originally added as a complex-ant, reduced the Ag� to Ag if used in high concentration. However, small amountswere found to give a more homogeneous film than if no triethanolamine was used.The deposition works best at room temperature—heating results in excessive pre-cipitation in solution. The iodopropanol usually contains some free iodide, theamount of which increases with age, and this can deleteriously affect the deposi-tion if present in too high quantities. The yellow AgI films exhibited sharp XRDpeaks (no line broadening), showing them to be a mixture of wurtzite and spha-lerite AgI. The films (other than very thin ones) scatter light moderately strongly,and transmission spectra were taken using an integrating sphere. Figure 7.4 showsthe spectrum of such a film. The strong absorption onset at 440–450 nm is due tothe direct bandgap of AgI (ca. 2.8 eV). AgI also has an indirect absorption atlonger wavelengths, and the decrease in transmission over this region is due partlyto this absorption and probably partly to scattered radiation not collected in the in-tegrating sphere.

The deposition occurs in parallel with homogeneous precipitation, suggest-ing that film formation is due to adhesion of crystals from the solution. This is sup-ported by SEM pictures that show scattered crystal formation, with gradual den-


sification of the crystal structure as the deposition proceeds. Figure 7.5 A showssuch a micrograph of a AgI film on glass. The average crystal size is a few hun-dred nanometers, and, as might be expected from such a morphology, the film ad-hesion is poor—the films can be wiped off with a tissue, although they will usu-ally stand up to cleaning in an ultrasonic bath. Adhesion is much better on

FIG. 7.4 Total transmission spectrum (measured with an integrating sphere) of a CDAgI film on glass.

FIG. 7.5 TEM micrographs of (A) a CD AgI film on glass and (B) a similar film onSnO2-coated glass.


SnO2-coated glass; the films are not removed by rubbing with a tissue. The filmmorphology is different than for the films on glass; the crystals are somewhatsmaller and tightly aggregated (Fig. 7.5B).

AgBr films could be made in the same manner using various bromo-alco-hols. AgBr is colorless but is usually slightly brownish due to photolytic forma-tion of small particles of metallic silver (this occurs much more strongly if depo-sition is carried out in room light, but formation of some brown coloration is stillnoticeable even if deposition is carried out in the dark). This coloration maskedthe optical absorption, which occurs mainly in the UV for AgBr. For bromides, tri-ethanolamine was not needed.

AgCl was not so readily deposited, and while occasionally some film for-mation did occur, it was nonhomogeneous, thin, and irreproducible. AgCl couldbe deposited by immersing a substrate in an aqueous solution of AgNO3 to whichNaCl solution was added, resulting in a cloudy precipitate/sol of AgCl. The con-centrations of both reactants were important: ca. 20 mM each, preferably with aslight excess of Ag. Much lower concentrations resulted in little deposition (atleast within a reasonable time), while even twice that concentration resulted in im-mediate coagulation of the AgCl and no film formation. What was unique aboutthis deposition was that on two occasions (among five or six experiments in total),visible film formation occurred virtually immediately on mixing the solutions.This is the only example known to the author where visible film formation occursin a rapid precipitation and contradicts the “conventional wisdom,” which other-wise seems to be valid, that the reaction leading to formation of the compoundmust be slow in order for appreciable film formation to occur.

More details on this work are given in Ref. 70.

7.4 EXTENSION OF CD TO OTHER MATERIALS?

To conclude this chapter, we look back at the earlier literature in hopes of widen-ing both the potential deposition methods and the materials that can be deposited.As well as oxides and related compounds, other anions are considered. The re-sulting compounds do not necessarily fall under the common heading of semi-conductors, but they are relevant in the hope of expanding the scope of chemicaldeposition.

Table 7.2 summarizes a range of homogeneous precipitation reactions. De-tails of all these reactions can be found in Ref. 2 (this book, in spite of its age, isrequired reading for anyone wishing to pursue this line; more recent books mayexist, but will probably not reduce its value). It should first be stressed that the ma-terial in this final section relates to precipitates rather than to films. However, withsome effort (in some cases only a little or none, as seen from the common film for-mation occurring in urea precipitations), it is reasonable to expect extension toform films of the same materials in at least some cases.


The list in Table 7.2 may appear incomplete to the modern chemist utilizingor studying chemical deposition; e.g., only thioacetamide is noted as a sulphidesource and selenides are not included. However, when we reflect that the vast bulkof the work carried out on CD concerned just sulphides, selenides and oxides, this“old” table might point the way to a major expansion of the CD technique, bothfor semiconductors and for other compounds. Further processing may be expectedto extend the types of material even further. For example, arsenates and phos-phates may be reducible in some cases to the better-known (to the semiconductorcommunity) arsenides and phosphides.

TABLE 7.2 Homogeneous Precipitation Reactions

Precipitant Reagent Elements precipitated

Hydroxide Urea Al, Ga, Th, Fe(III), Sn, ZrAcetamide TiHexamethylenetetramine ThMetal chelate � H2O2 Fe(III)

Phosphate Triethyl phosphate Zr, HfTrimethyl phosphate ZrMetaphosphoric acid Zr

Oxalate Dimethyl oxalate Th, Ca, Am, Ac, rare earthsDiethyl oxalate Mg, Zn, CaUrea and an oxalate Ca

Sulphate Dimethyl sulphate Ba, Ca, Sr, PbSulphamic acid Ba, Pb, RaPotassium methyl sulphate BaAmmonium persulphate BaMetal chelate � persulphate Ba

Sulphide Thioacetamide Pb, Sb, Bi, Mo, Cu, As, Cd,Sn, Hg, Mn

Iodate Iodine � chlorate Th, ZrPeriodate � ethylene diacetate Th, Fe(III)

(or �-hydroxyethyl acetate)Carbonate Trichloroacetate Rare earths, Ba, RaChromate Urea � dichromate Ba, Ra

Potassium cyanate � dichromate Ba, RaCr(III) � bromate Pb

Chloride Ag–ammonia complex � chloride � Aghydroxyethylacetate

Arsenate Arsenite � nitric acid ZrFluoride Fluoboric acid La

Source: Modified from: I. M. Kolthoff and P. J. Elving, eds. Treatise on Analytical Chemistry, Part 1,Vol. 1. New York: Interscience Encyclopedia, 1959, p 741.


REFERENCES1. L Gordon. Anal. Chem. 24:459, 1952.2. L Gordon, ML Salutsky, HH Willard. Precipitation from Homogeneous Solutions.

New York: Wiley, 1959.3. HH Willard, NK Tang. Ind. Eng. Chem. Anal. Ed. 9:357, 1937.4. HH Willard, HC Fogg. J. Am. Chem. Soc. 59:40, 1937.4a. HH Willard, L Gordon. Anal. Chem. 25:170, 1953.5. TP Niesen, MR De Guire. J. Electroceram. 6:169, 2001.6. RN Bhattacharya, P Pramanik. J. Electrochem. Soc. 129:1642, 1982.7. RL Call, NK Jaber, K Seshan, JR Whyte Jr. Sol. Energy Mater. 2:373, 1980.8. M Najdoski, I Grozdanov, B Minceva. J. Mater. Chem. 6:761, 1996.9. M Ortega, G Santana, A Morales-Acevedo. Solid State Electron. 44:1765, 2000.

10. M Ocampo, AM Fernandez, PJ Sebastian. Semicond. Sci. Technol. 8:750, 1993.11. M Ocampo, PJ Sebastian, J Campos. Phys. Status Solidi (a) 143:K29, 1994.12. AJ Varkey, AF Fort. Thin Solid Films 239:211, 1994.13a. FC Eze. J. Phys. D: Appl. Phys. 32:533, 1999.13. AJ Varkey, AF Fort. Sol. Energy Mater. Sol. Cells 31:277, 1993.14. B Pejova, A Isahi, M Najdoski, I Grozdanov. Mater. Res. Bull. 36:161, 2001.15. I Grozdanov. Mater. Lett. 19:281, 1994.16. M Ristov, G Sinadinovski, I Grozdanov. Thin Solid Films 123:63, 1985.17. RP Goyal, D Raviendra, BRK Gupta. Phys. Status Solidi (a) 87:79, 1985.18. M Izaki. Electrochem. Solid State Lett. 1:215, 1998.19. KO Velthaus, J Kessler, M Ruckh, D Hariskos, D Schmid, HW Schock. In: 11th

ECPV Solar Energy Conf., Montreux, Switzerland, 1992, p 842.20. PC Rieke, BD Marsh, LL Wood, BJ Tarasevich, J Liu, L Song, GE Fryxell. Langmuir

11:318, 1995.21. BJ Tarasevich, PC Rieke, J Liu. Chem. Mat. 8:292, 1996.22. PC Rieke, R Wiecek, BD Marsh, LL Wood, J Liu, L Song, GE Fryxell, BJ Tarase-

vich. Langmuir 12:4266, 1996.23. QW Chen, YT Qian, H Qian, ZY Chen, WB Wu, YH Zhang. Mater. Res. Bull.

30:443, 1995.24. QW Chen, YT Qian, ZY Chen, Y Xie, GE Zhou, YH Zhang. Mater. Lett. 24:85, 1995.25. QW Chen, XG Li, YT Qian, YH Zhang. Mater. Lett. 31:247, 1997.26. B Pejova, M Najdoski, I Grozdanov, A Isahi. J. Mater. Sci.-Mater. Electron. 11:405,

2000.27. M Izaki, O Shinoura. Adv. Mater. 13:142, 2001.28. P Ruëtschi, B Cahan. J. Electrochem. Soc. 104:406, 1957.29. W Mindt. J. Electrochem. Soc. 117:615, 1970.30. RN Bhattacharya, P Pramanik. Bull. Mater. Sci. 2:287, 1980.31. S Gorer, A Albu-Yaron, G Hodes. Chem. Mater. 7:1243, 1995.32. P Pramanik, S Bhattacharya. J. Mater. Sci. Lett. 6:1105, 1987.33. KL Chopra, RC Kainthla, DK Pandya, AP Thakoor. In: Physics of Thin Films, Vol.

12, Academic Press, New York and London, 1982, p 167.34. W Mindt J. Electrochem. Soc. 118:93, 1971.35. E Beutel, A Kutzelnigg. Monats. 58:295, 1931.


36. P Pramanik, S Bhattacharya. J. Electrochem. Soc. 137:3869, 1990.37. AJ Varkey, AF Fort. Thin Solid Films 235:47, 1993.38. B Pejova, T Kocareva, M Najdoski, I Grozdanov. Appl. Surf. Sci. 165:271, 2000.39. H Nagayama, H Honda, H Kawahara. J. Electrochem. Soc. 138:2013, 1988.40. AJ Varkey, AF Fort. Sol. Energy Mater Sol. Cell 29:253, 1993.41. T Kocareva, I Grozdanov, B Pejova. Mater. Lett. 47:319, 2001.42. D Raviendra, JK Sharma. J. Phys. Chem. Solids 46:945, 1985.43. QW Chen, YT Qian, ZY Chen, GE Zhou, YH Zhang. Thin Solid Films 264:25, 1995.44. D Hariskos, R Heberholz, M Ruckh, U Ruhle, R Shäffler, HW Schock. In: 13th

ECPV Solar Energy Conf., Nice, France, 1995, p 1995.45. S Supothina, MR De Guire. Thin Solid Films 371:1, 2000.46. K Tsukuma, T Akiyama, H Imai. J. Non-Cryst. Sol. 210:48, 1997.47. BC Bunker, PC Rieke, BJ Tarasevich, AA Campbell, GE Fryxell, GL Graff, L Song,

J Liu, JW Virden, GL McVay. Science 264:48, 1994.48. H Shin, RJ Collins, MR Deguire, AH Heuer, CN Sukenik. J. Mater. Res. 10:692,

1995.49. RJ Collins, H Shin, MR DeGuire, AH Heuer, CN Sukenik. Appl. Phys. Lett. 69:860,

1996.50. QW Chen, YT Qian, ZY Chen, WB Wu, ZW Chen, GE Zhou, YH Zhang. Appl. Phys.

Lett. 66:1608, 1995.50a. S Deki, Y Aoi, Y Miyake, A Gotoh and A Kajinami. Mater. Res. Bull. 31:1399, 1996.51. M Agarwal, MR DeGuire, AH Heuer. Appl. Phys. Lett. 71:891, 1997.52. D Raviendra, JK Sharma. J. Appl. Phys. 58:838, 1985.53. T Saeed, P O’Brien. Thin Solid Films 271:35, 1995.54. P O’Brien, T Saeed, J Knowles. J. Mater. Chem. 6:1135, 1996.55. J McAleese, P O’Brien. Mater. Res. Soc. Symp. Proc. Vol. 485. 1998, p 255.56. QW Chen, YT Qian, ZY Chen, YH Zhang. Mater. Lett. 22:93, 1995.57. A Ennaoui, M Weber, R Scheer, HJ Lewerenz. Sol. Energy Mater. Sol. Cells 54:277,

1998.58. AJ Varkey. Int. J. Mater. Prod. Technol. 10:94, 1995.59. M Izaki, T Omi. J. Electrochem. Soc. 144:L3, 1997.60. M Izaki, J Katayama. J. Electrochem. Soc. 147:210, 2000.61. M Agarwal, MR De Guire, AH Heuer. J. Am. Ceram. Soc. 80:2967, 1997.62. G Contreras-Puente, O Vigil, M Ortega-Lopez, A Morales-Acevedo, J Vidal, ML Al-

bor-Aguilera. Thin Solid Films 361:378, 2000.63. VP Tolstoy. Thin Solid Films 307:10, 1997.64. GA Kitaev, GM Fofanoy. Zh. Prikl. Khim. 43:1694, 1970.65. B Pejova, I Grozdanov. Appl. Surf. Sci. 177:152, 2001.66. M Perakh, A Peled, Z Feit. Thin Solid Films 50:273, 1978.67. M Perakh, A Peled. Thin Solid Films 50:283, 1978.68. M Perakh, A Peled. Thin Solid Films 50:293, 1978.69. A Peled. Philosoph. Mag. B 53:171, 1986.70. G Hodes, G Calzaferri. Adv. Funct. Mater. in press.


8Ternary Semiconductors

8.1 GENERAL CONSIDERATIONS FORPRECIPITATION AND CHEMICALDEPOSITION OF TERNARY COMPOUNDS

Mixed compositions are of interest mainly because they allow tuning of the semi-conductor properties (most commonly bandgap and, therefore, spectral sensitiv-ity). This is useful for various device applications. Photoconductive detectors,where a certain spectral sensitivity range is desired, is probably the main applica-tion that drove many studies on CD of ternary semiconductors.

Mixed metal chalcogenides have been deposited by CD. According to sim-ple fundamental considerations, the deposition should proceed according to thesolubility products of the two separate metal chalcogenides; the one with a smallerKsp should precipitate first, and only after, when the concentration of free (first)metal ion was low enough, would the other chalcogenide precipitate, assuming asufficient supply of chalcogenide ions. If the difference in Ksp was large, then thesolution would be almost entirely depleted of the low-Ksp chalcogenide beforeprecipitation of the second would start. Note that this discussion relates to CDwhere deposition is slow; for rapid precipitation, kinetic factors might be more im-portant, and differences in concentrations of the two cations are likely to play amore dominant role.


This picture is in many (probably most) cases oversimplified. There areother factors that may become important in such mixed products. The literature isfull of examples of precipitation of mixtures/compounds, and we will considersome of these, keeping in mind that they refer to precipitation and not CD. One iscoadsorption of one cation on the precipitated compound of the other. A well-known example of this phenomenon is adsorption of Cr3� on ferric hydroxide pre-cipitated from solutions containing Fe3� and small concentrations of Cr3�. Thechromium is precipitated together with the Fe(OH)3 at a pH where chromium hy-droxide itself would not readily precipitate [1]. This is due to strong adsorption ofthe “impurity” cation (Cr3�) on the precipitating Fe(OH)3. The ability of someprecipitates to either cause coprecipitation of other metal salts that would, bythemselves, not precipitate under the same conditions—an effect that was knownas induced precipitation—or often even incorporate the “soluble” metal salt if thefreshly precipitated insoluble salt was exposed to a solution containing the moresoluble one—has been known for a long time; many examples from the old liter-ature are given in a review by Kolthoff and Moltzau [2].

Another, and on the face of it, rather different example, is the coprecipita-tion of solid solution compounds, such as CuInS2 and CuInSe2—semiconductorsof particular interest due mainly to their applicability for photovoltaic cells. It wasshown, by X-ray diffraction, that the precipitate resulting from reaction betweenH2S and an aqueous solution containing both Cu� and In3� ions was, at least inpart (depending on the concentrations of the cations), single-phase CuInS2 [3].Two factors were found to be necessary for this compound formation: (1) the pres-ence of sulphide on the surface of the initially precipitated colloidal solid metalsulphide and (2) one of the cations being acidic and the other basic. The monova-lent Cu� cation is relatively basic, while the trivalent In3� cation is relativelyacidic. It is not clear what the physical reason is for this latter requirement. A dif-ference in practice between acidic and basic cations is that, in an aqueous solutionof both cations, the acidic cation is more likely to be in the form of some hydroxyspecies (not to be confused with hydrated cations), while the basic cation is morelikely to exist as the free cation.

In a subsequent study of AgMGaMS precipitation by the same group, amechanism for solid solution formation was proposed [4]. Colloidal Ag2S, withits lower solubility product, formed initially. The surface of the colloid adsorbedboth sulphide and Ga (or Ga-hydroxy; see earlier) ions to form a gallium sulphidelayer. This then would be similar to induced precipitation. It was suggested thatthis GaMS then diffused into the Ag2S, where solid solution formation occurred.In view of the high mobility of Ag� ions, it seems more likely that Ag� diffusedoutward rather than the Ga-species diffusing inward. Of course, precipitation is ahighly nonequilibrium process, while CD, depending on which mechanism is op-erative, is closer to an equilibrium process (an ion-by-ion deposition occurs closeto equilibrium, while the initial hydroxide formation in a hydroxide mechanism is


more like a precipitation reaction; the subsequent chalcogenide exchange is againcloser to equilibrium). Thus, extension of precipitation reactions to CD processesmay be useful, but the comparison should be made cautiously.

Various possibilities for precipitation of two cations, M1 and M2, by sul-phide are shown in Figure 8.1. The case where both metal sulphides have similarsolubility products (Cd,PbS would be an example of this) is shown in the left-handprocess, where a particle of M1M2S is formed (to simplify things, no informationon stoichiometry is implied here). This particle may be either a mixed phase or asolid solution, depending on the miscibility of the two sulphides and the kineticsof the precipitation. The middle process shows the case for two sulphides with dif-ferent solubility products. The metal sulphide with the lower-solubility product(assumed to be M1) precipitates as a separate phase, which adsorbs both sulphideions and M2 ions (the latter either onto adsorbed S or onto lattice S). Eventually ashell of M2S (probably containing M1S) will form. Depending on the drivingforces involved, this core-shell structure may remain in that state, or diffusion intoa mixed two-phase particle or single-phase solid solution may occur. A third pos-sibility, shown on the right-hand side, is that only M1S is formed. If the M1 is suf-ficiently depleted by precipitation of M1S, and if sulphide formation continues,then M2S will eventually precipitate.

These processes have been shown for free metal ions. However, if a clustermechanism based on metal hydroxide colloids is involved, they are equally appli-cable to the formation of the solid hydroxide species. The degree of conversion of

FIG. 8.1 Scheme of various possibilities for coprecipitation of two metal (M1 and M2)sulphides.


the hydroxides to sulphide will then depend on the differences between the solu-bility products of the hydroxides and sulphides as well as the hydroxide concen-tration in solution (pH and temperature). If a complex-decomposition mechanismis involved, then the most important factor is probably the differences between thestrengths of the metal–sulphur (keeping sulphide as our example) bonds, and thisis similar to the differences in solubility products of the sulphides.

These processes have been described for rapid precipitation reactions. How-ever, they should also be valid in general for slow precipitation—i.e., for CD—with possible differences due to the very different kinetics involved. Thus, if freesulphide is involved, since it is always present in very low concentration, thelower-solubility product metal sulphide is more likely to deposit first, comparedto rapid precipitation. Solid-state diffusion processes have much more time to oc-cur in CD (although they may occur in rapid precipitation after the precipitationitself), increasing the probability of solid solution formation.

Probably the most basic question to be asked when depositing ternaries is: Isthe as-deposited material a true single-phase solid solution? A look at Table 2.3shows that this is often not the case. In some studies, clear structural characteriza-tion (usually XRD) shows that the film is (at least predominantly) a single-phasesolid solution. Others state that the product is not a solid solution or contains a largecomponent of other phases. However, some studies claim that the films are solidsolution, without presenting clear evidence. In most cases, these are older studies,based on XRD spectra which, by today’s standards, are not clear. The XRD spec-trometers of today are a lot better than those of not so long ago. Additionally, nowa-days it is recognized that nanocrystalline films (as usually deposited by CD) oftenrequire more care in sampling, to avoid the common danger of incorrectly pro-nouncing them “amorphous” or “poorly crystalline.” Films that are not a true solidsolution as deposited may often be converted to a solid solution by annealing. Sim-ilarly, solid solution formation may sometimes occur if two separate films are de-posited, one on top of the other, and then annealed to effect interdiffusion. Whilethe original intention of this chapter was to confine the contents to genuine solidsolutions (or at least those that might be solid solutions), in some cases ternaryfilms that clearly are not solid solutions are included, and this is then made clear.

Another consideration is whether the deposited film is homogeneousthroughout its thickness. If the composition is a function of the relative solubilityproducts of the individual binary compounds, then the metal ion that has thelower-solubility product with the anion will deposit preferentially at first, but, dueto depletion, it may become lower in concentration in the film as deposition pro-ceeds. Thus for complete characterization of these films, compositional analysisshould ideally be made as a function of spatial position in all three dimensions;this is rarely carried out in practice.

The semiconductors described in this chapter are divided into two types:those composed of two different metal cations (most of the studies) and those withtwo different anions. In their 1982 review [5], Chopra et al. give a list of eight dif-


ferent ternary compounds that had been deposited by CD at that time. The workon some of those compounds has apparently not been published, other than in the-ses, and details of their preparation and characterization are not given; thereforethose studies are not discussed here. These include the sulphides of Cd,Zn andCd,Hg and the selenides of Cd,Hg, Cd,Pb, and Pb,Hg. Additionally, two papers onbilayer formation are treated separately in Section 8.4.

8.2 MIXED METAL COMPOUNDS

8.2.1 (Cd,Zn)O

(Cd,Zn)O films were deposited on glass at 45°C from solutions of Cd and Zn chlo-rides to which ammonia and then H2O2 were added (the purpose of the H2O2 wasnot given); they were then annealed at a final temperature of 500°C [6]. (No de-scription of the as-deposited films was given; they were presumably mixed hy-droxides.)

Optical spectroscopy showed that the optical bandgap shifted strongly to thered with increasing Cd concentration (at fixed Zn concentration) for small Cd:Znratios in the deposition bath, but it was affected only to a small extent by varia-tions in this ratio when the Cd:Zn ratio was greater than ca. unity. From this it canbe inferred that Cd(OH)2 was preferentially deposited, even though the solubilityproduct of Zn(OH)2 is lower (by ca. 50 times). This could be explained by thegreater strength of the Zn-ammine complex compared to the corresponding Cdcomplex (two orders of magnitude higher), resulting in a hundred times lowerfree-Zn2� concentration compared to Cd (for the same total concentration ofeach). This more than offsets the lower-solubility product of the Zn(OH)2, al-though the difference is not large, resulting in preferential deposition of Cd(OH)2.This provides a good example of the need to consider all the relevant parameterswhen trying to understand the specifics of the depositions.

The resistivity of pure CdO was 3 � 10�3 �-cm (CdO is normally a de-generate n-type semiconductor), which increased approximately linearly (on asemilog scale) with increasing solution Zn content up to 107 �-cm (at 60% Zn)and then tailed off to a value of ca. 108 �-cm for very Zn-rich films.

8.2.2 (Cd,Zn)S

By far the greatest interest and effort in the CD of ternary semiconductors has beenfocused on cadmium zinc sulphide (Cd,Zn)S. This interest has been driven by theexpected improvement in performance of thin-film photovoltaic cells (CdTe- andCIS (CuInSe2-based cells) using (Cd,Zn)S rather than the presently used CdS.This expectation arises mainly from the increased bandgap of the Zn-containingsolid solution, resulting in increased transparency to shorter wavelengths of light(see Sec. 9.1.4.5 for more details). Another consideration for heterojunction for-


mation is the decrease in electron affinity of the semiconductor with increase inZn content (ZnS has a smaller electron affinity than CdS). The electron affinity ofa semiconductor is a measure of the position of the conduction band with respectto the vacuum energy level; a lower value means a higher conduction band. Thusthe alignment of the conduction band of the (Cd,Zn)S with that of the secondsemiconductor can be controlled to a large extent by varying the film composition.

In spite of the overall chemical similarity of Cd and Zn, however, it has notproven simple to deposit true solid solutions of the sulphides. There are a numberof reasons for this, some of which have been treated in detail in Chapter 4 in thediscussion on ZnS deposition. We sum them up here.

While CdS is less soluble than ZnS, Cd(OH)2 is more soluble than Zn(OH)2.For this reason, ZnS is more difficult to deposit than CdS, since Zn(OH)2 tends toform instead of, or together with, ZnS. (Although the solubility product of ZnS islower than that of Zn(OH)2, the concentration of hydroxide in any typical aque-ous solution will be much higher than that of sulphide). In an alkaline solution (themost common medium for CD), CdS deposition will be preferred over ZnS.

The concentration of free Cd2� should be much lower than that of Zn2� inorder for ZnS to deposit according to simple solubility product considerations.However, the strength of complexation of most ligands is comparable for both Cdand Zn (ammonia and hydroxide give stronger complexes with Zn). We know ofno ligand that will complex Cd enough to bring its free-ion concentration in solu-tion the orders of magnitude lower than that of Zn in the same solution that is re-quired (cyanide is maybe the closest to this ideal, but the difference is still notenough, and cyanide is such a strong complexant that deposition might be rela-tively difficult from solutions containing it in large amounts). Because of thelower solubility of Zn(OH)2, it should be possible to adjust the complexant con-centration so that Zn(OH)2 is present in solution but Cd(OH)2 is not, ignoring thepossibility of induced coprecipitation. CdS would therefore be formed by the(usually slower) ion-by-ion mechanism, while ZnS might be formed by the (usu-ally faster) hydroxide cluster mechanism. This is probably not as ideal as it maysound, both because Zn(OH)2 does not readily methasize to ZnS (due to the muchhigher hydroxide than sulphide concentration in the solution) and because, if it didoccur, it is more likely that separate phases would be formed.

One point in favor of a single-phase solid solution deposition is that CdS andZnS do readily form solid solutions in general. Thus, if the two sulphides can besimultaneously deposited, there is a good chance that they will form a solid solu-tion if the temperature is high enough. Put another way, annealing of a well-mixedtwo-phase mixture of CdS and ZnS will form a solid solution if the temperature ishigh enough. For finely divided precipitates (as normally occurs in CD), this tem-perature is expected to be relatively low.

There are a number of reports on CdMZnMS deposition where the CD filmwas either clearly shown to be mixture of phases [7] or there was insufficient ev-


idence to support solid solution formation [8,9]. The first well-characterized de-position of a true (Cd,Zn)S alloy film was described by Padam et al. [10]. Cd andZn acetates were used in various ratios complexed with ammonia and tri-ethanolamine and thiourea at 90–95°C to deposit (Cd,Zn)S over the completecomposition range onto glass substrates. Interestingly, the Zn will be more heav-ily complexed than the Cd in this solution, which shows that the mechanism of de-position is not one based solely on solubility products of the sulphides. In fact,from the crystal size measurements of a similar deposition described in Ref. 11(see later), it is possible that the deposition mechanism is different for the twocations (the small crystals of pure ZnS and the larger ones of pure CdS suggest acluster mechanism and an ion-by-ion mechanism, respectively).

The films were characterized by a variety of techniques. Elemental analysis(EDS) showed that the Zn:Cd ratio in the film was almost equal (slightly less) thanthat in the deposition solution. X-ray diffraction and ED were used for phase andcompositional analyses. All the compositions up to 80% Zn were wurtzite struc-ture, while pure ZnS was sphalerite. Interestingly, while most of the films gavering ED patterns showing nonoriented growth, some showed a degree of orienta-tion, in spite of the glass substrate.

The bandgap, calculated from optical absorption spectroscopy, varied al-most linearly with composition between that of CdS (2.4 eV) and ZnS (3.6 eV),providing further evidence for solid solution formation.

The films were all n-type (hot probe) with resisitivity � that varied linearly(on a log � scale) from 109 �-cm (CdS) to 1012 �-cm (ZnS). Doping by In (asInCl3 in the deposition solution) reduced �; e.g., for a Cd0.8Zn0.2S film, � droppedlinearly (on a log � scale) with In content from ca. 1010 �-cm (undoped) to ca. 105

�-cm (1.5% In—the In ratio in the film was similar to that in the solution). At highIn ratios, � increased, explained by a decrease in mobility due to scattering by In.Annealing in H2 at 200°C also decreased �. For example, a Cd0.8Zn0.2S:1.5% Infilm showed a minimum value for � of ca. 10 �-cm, presumably due to loss of S.

This same method was more recently repeated with very similar results[11]. It was additionally found that the films were strongly textured (only oneXRD peak—either (0001) wurtzite or (111) sphalerite), although this texturewas lost if a subsequent layer was deposited to produce thicker films. The crys-tal size (measured from XRD peak width) varied from 20 nm (CdS) to 9 nm(ZnS). The bandgap varied between the same limits as found in the previousstudy, but changed more rapidly for high Zn content. The resistivity of the filmsvaried (linearly on a log � scale) from 109 �-cm (CdS) to 1014 �-cm (ZnS)—the latter higher than the value measured by Padam et al. Boron doping (addi-tion of boric acid to the solution) decreased the resistivity of CdS by three or-ders of magnitude.

Using ammonia-complexed metal iodides and thiourea at pH 10, films wereformed whose properties depended on the temperature–time regime of the depo-


sition solution [12]. If the reagents were mixed at room temperature and deposi-tion occurred while heating the solution to 80°C, only CdS was deposited fromvarious mixtures of Cd and Zn iodides, even where the Zn was in large excess. If,however, the reagents were mixed at a higher temperature (60°C), followed byheating the bath to 80°C during deposition, Zn was incorporated into the films, al-though the Zn content in the films was much lower than in the solutions; the Zncontent increased slowly up to ca. 80% Zn in solution and then rapidly at higherZn concentrations. This was seen by both optical spectra and XRD, the latter inparticular supporting solid solution formation. A reason for this dependence ontemperature programming was not suggested. A possible clue may be obtainedfrom a consideration of the temperatures at which solid Cd(OH)2 and Zn(OH)2 areformed (from thermodynamic calculations based on the equilibrium constants ofthe Cd- and Zn-ammine complexes, the solubility products of the hydroxides, anddependence of hydroxide concentration in water on temperature). At room tem-perature, no solid hydroxide phase is calculated to be present in the various Cd/Znsolutions. At approximately 40°C, hydroxide will form (slightly lower tempera-ture for Cd, slightly higher for Zn, although the difference is not large, and, inview of the approximations used in these calculations, as well as kinetic factors, itis not certain that Cd(OH)2 will, indeed, form first). Also, it has been shown thatCd(OH)2 can form on the substrate in some cases before it forms in solution (dis-cussed in Chap. 3), but the equivalent experiment has not been done for Zn. How-ever, if we assume that Cd(OH)2 will form before Zn(OH)2 as the temperature israised, then this might explain, at least in part, the formation of only CdS. For thecase where deposition was started at 60°C, then both hydroxides are present to be-gin with and it is more likely that Zn will be incorporated into the final films. Thisreasoning is based on the expectation that cluster deposition will be much fasterthan ion-by-ion deposition.

The resistivity of the films decreased with increase in Zn content (from 1010

�-cm for CdS (a very high value for CdS) to ca. 106 �-cm for 90% (solution con-centration) Zn and then increased to ca. 109 �-cm for pure ZnS. No explanationfor this effect was given. The films were photoconductive, with the resistivity de-creasing in a somewhat sporadic manner as a function of composition, up to amaximum dark:light ratio of 5 � 103 for the 90% Zn films.

Similar conditions, but at a lower pH of 8.4, were also used by the samegroup [13]. The ammonia concentration was reported to be important in formingthe solid solutions, although this concentration was not given. Only Zn-rich solu-tions were described in this study (between 80 and 99% Zn in solution). As for theprevious study at higher pH, the Zn concentrations in the films were quite differ-ent from those in the solution (except for very Zn-rich solutions); the films werericher in Cd up to at least 92% Zn in solution, and then the Zn concentration in thefilms increased rapidly with further increase in solution Zn concentration up to99% Zn, from which deposited films with very little Cd.


X-ray diffraction showed well-defined peaks that shifted in position withchange in composition, while the optical absorption spectra gave values of Eg thatalso varied gradually with composition. The well-defined shift in XRD peaks and,to a somewhat lesser extent, the gradual change in estimated bandgap with com-position provide good evidence for true solid solution formation.

Unlike other studies as well as the higher-pH studies by the same group, theelectrical resistivity did not vary much with composition, being ca. 5 � 107 �-cm.An overview of the variation of resistivity of some of these (Cd,Zn)S films withcomposition is given in Figure 8.2. There is quite a large variation, both in resis-tivity values and in their compositional dependence.

The question must be asked: Why are solid solutions formed in some casesand not in others? A common denominator in the successful films and their dif-ference from the unsuccessful ones (success being defined as formation of a solidsolution) is the higher temperature used in the former (80–95°C). Higher temper-ature will facilitate intermixing of the codeposited CdS and ZnS.

8.2.3 (Cd,Zn)Se

Two selenosulphate baths have been described for (Cd,Zn)Se. In the first [14],solid solution formation was claimed, at least for annealed (300°C) samples, al-

FIG. 8.2 Resistivity data for some (Cd,Zn)S films. The data from the two Yamaguchipapers were modified to show resistivity values as a function of approximate film compo-sition, rather than solution composition as given in the original papers.


though XRD results of the (annealed) samples were difficult to understand (the“a” lattice parameter increased as Zn concentration increased). The effects ofcomposition and annealing on the bandgaps and electrical conductivity were de-scribed. In the other [15], solid solution formation over part of the compositionrange was claimed. However, the XRD spectra of the films (relatively sharp peakscharacteristic of ZnO), together with the visual description of the pure ZnSe films(white with a slight greyish tinge), suggest mixed-phase formation containingZnO.

8.2.4 (Cd,Hg)S

Solid mixtures of CdS and HgS have been shown to form solid solutions aftertreatment with certain solutions, such as concentrated ammonium sulphide [16].This may be due partly to the very similar ionic radii of the two cations and(maybe more important) the ability of Hg to diffuse readily in solids. Therefore itis probable that solid solutions can readily form in this system.

(Cd,Hg)S was deposited from a solution containing CdCl2 and HgCl2 com-plexed with a low concentration of cyanide (15 mM CN� to 50 mM CdCl2—theCd concentration was fixed and the Hg was varied), thiourea, and ammonia to-gether with KOH at 80–85°C onto Ti substrates [17].

The mole fraction of Hg in the films was ca. four times its concentration inthe solution. This was expected based on the lower-solubility product of HgScompared to CdS. The maximum mole fraction of Hg in the films (compositionmeasured by atomic absorption spectroscopy) was 0.18; attempts to increase thisvalue by adding more HgCl2 resulted in rapid precipitation in the solution and lit-tle film formation. Increased Hg concentration in the films could probably be ob-tained by optimization of the conditions, e.g., by reducing the Cd concentrationand/or by using a specific complex for Hg, such as iodide. The bandgap of thefilms (annealed at 320°C for 3 hr in air), measured by photoelectrochemical pho-tocurrent spectroscopy, decreased with increasing Hg content down to 1.8 eV for0.18 mole fraction Hg. The shape of these spectra suggested that the (annealed)films were solid solutions, although no structural characterization was made. Themain purpose for making these films seems to have been to study their photoelec-trochemical properties, which are described in Chapter 9.

Triethanolamine was also used as a complexant to deposit these films fromthiourea baths [18]. As with the previous study, there was a maximum Hg contentin the bath (0.05 mole fraction—absolute concentrations were not given), whichled to a 0.18 Hg mole fraction in the films, above which, although films wereformed, the Hg content decreased, also explained by rapid precipitation of HgS inthe solution. X-ray diffraction showed the formation of a single phase, up to a Hgcontent (in the bath) of 0.15, and two-phase formation at higher concentrations.The optical bandgap dropped from 2.4 eV (pure CdS) to 1.76 eV (0.05 Hg in bath,


0.18 in film) and then slowly increased again with increasing Hg in the bath to ca.1.9 eV.

8.2.5 (Cd,Pb)S

Solid solutions of Cd and Pb sulphides have been popular in ternary CD. The crys-tal structures of the individual sulphides are different: PbS crystallizes in the rock-salt structure, while CdS forms tetrahedrally bonded sphalerite or wurtzite struc-ture (a rocksalt form of CdS exists but normally only at high pressures). Thissuggest that the solubility range of the alloys will be limited.

Acetates of Cd and Pb were mixed with ammonia and thiourea and films de-posited on glass at room temperature [19]. The concentration of Cd in the filmswas a little higher than in the solution. The lattice constant increased with increasein Cd but was greater than that of PbS (which has a slightly larger lattice constantthan CdS) at all levels measured. This is not typical of solid solution formation,although it does imply a single phase. It was suggested that lattice expansion oc-curred due to PbS entering as an interstitial into the CdS. It seems more probablethat, if interstitial expansion is the correct explanation, Cd2�, which is the lower-concentration component, will go into interstitial sites in the PbS. The morphol-ogy of these layers was investigated as a function of composition [20].

In a similar study, (Cd,Pb)S was deposited on Ti at ca. 75°C [17]. The ratioof Cd to Pb in the films was found (by atomic absorption spectroscopy) to be verysimilar to that in the solution, a consequence of the similar values for the solubil-ity product of the two sulphides. Since the main purpose for investigating thesefilms was to study their photoelectrochemical properties (see Chap. 9 for details),little no characterization, other than compositional and photoelectrochemical, wasmade. Photoelectrochemical spectroscopy (of films annealed at 460°C in air for 4hr) showed a decrease in bandgap with increasing Pb content down to 1.6–1.7 eVfor Cd0.82Pb0.18S, although very nonlinearly—a strong drop in bandgap, of0.5–0.6 eV, occurred between 0.1 and 0.18 mole fraction Pb.

Films deposited from mixed Cd/Pb solutions complexed with ammonia (forCd) and hydroxide (for Pb), both in a minimum amount to effect dissolution, at pHvalues between 10 and 13 and deposition temperatures between 60 and 80°C wereconcluded, from consideration of the XRD, TEM (which showed two differentcrystal sizes), and optical spectra, to be mixtures of the two sulphides rather thana solid solution [21]. A study of these films for solar-selective surfaces was car-ried out.

A study of variation of the composition of (Cd,Pb)S films as a function ofthe free-[Cd2�]:[Pb2�] ratio (i.e., the uncomplexed metal ions, which could be cal-culated from the concentrations of total metal ions, concentration of complex, andthe respective stability constants of the metal complexes) showed a linear increasein Cd content of the films up to a ratio value of 10, followed by a sudden decrease


in Cd content above this value (the details of the solution composition were notgiven) [22]. The maximum amount of Cd that could be incorporated as a solid so-lution with PbS was 15%. It was suggested that for [Cd2�] greater than this, therate of CdS formation was so great that it became more favorable for the CdS toform a separate phase than a solid solution.

The crystallographic texture of the films was dependent on the Cd content.Up to 3 at.%, the films were (111) textured, while for higher Cd concentrationsthey became (200) textured. The crystal size (measured from electron mi-croscopy) was of the order of some hundreds of nanometers (somewhat smallerfor larger Cd content) but increased again to ca. 1 �m for maximum Cd contentjust before phase separation.

The resistivity of the films increased from 106 �-cm for very low Cd con-tent to a maximum of ca. 108 �-cm at 6% Cd and then slowly dropped again withincreasing Cd. This maximum correlates with the minimum crystal size, suggest-ing a dominant role of grain boundaries in the conduction mechanism. The spec-tral response of the photoconductivity blue-shifted with increase Cd content up toa peak response at 1.35 �m for 8.4% Cd.

The presence of cyanamides of Cd and Pb in films of (Cd,Pb)S was con-firmed by thermal desorption mass spectrometry [23]. Cyanamide (H2CN2) is aproduct of the decomposition of thiourea and forms sparingly soluble metal salts.The metal cyanamide content of the film varied from ca. 5% up to ca. 20% (byweight). The presence of the cyanamides decreased the intensity of the XRD re-flections, presumably due to poorer crystallization of the sulphides. Interestingly,the photosensitivity of the films increased with higher metal cyanamide content,although whether this was due specifically to the presence of the cyanamide or toits effect on the crystal growth was not known.

The same group also deposited (Cd,Pb)S using a flow system [24]. In thiscase, metal cyanamides were not detected by XRD, presumably because theflow system removed the cyanamide. The rate of flow affected the crystal size:Larger flow rate resulted in finer-grained deposits. Elemental analysis and XRDshowed the incorporation of Cd in the films, again up to ca. 10%, as a solid solution.

8.2.6 (Hg,Pb)S

Films of (Hg,Pb)S were grown on glass at 30°C from a solution of PbAc2, HgCl2,thiourea, and NaOH at pH � 10 [25,26]. It was noted that the pH had to be criti-cally controlled to obtain good-quality films. Additionally, the order of mixing ofthe solutions was unusual (it was not stated if this was critical or not) in that thePbAc2 and thiourea were first mixed, the NaOH was then added until a light browncolor appeared in the solution, and only then was the HgCl2 solution added, fol-lowed by adjustment of pH to 10.


For films deposited at temperatures up to 45°C, ED showed the presence ofa single phase, with increasing lattice spacing with increased Hg content up to 4%Hg. Above this concentration, aggregates of �-HgS (metacinnabar) were found inthe deposit, which originated from the colloidal solution. While not clearly de-fined, the transmission spectra shifted to the blue (increasing bandgap) with in-crease in Hg concentration. This increase in bandgap, together with the ED data,suggests that the films are alloys of PbS with �-HgS (cinnabar, bandgap ca. 2 eV).For deposition temperatures greater than 45°C, the lattice parameters decreasedwith increasing Hg, and the �-HgS phase was formed (the lattice parameters of �-HgS are ca. 1.5% smaller than those of PbS).

If FeCl3 was added to the HgCl2 solution, the film properties, even for thosedeposited at low temperatures, were similar to those grown without FeCl3 athigher temperatures. Single-phase films were obtained up to 50 at.% Hg with the�-HgS phase, and the lattice spacing decreased with increasing Hg content. Thetransmission spectra, while again mostly not well defined, shifted to the red withaddition of Hg. The bandgap of �-HgS was taken to be ca. 0.1 eV in these stud-ies; other values of ca. 0.5 eV have also been measured for this phase. The FeCl3was believed to stabilize the �-phase.

A specific study of the optical and electrical properties of these Pb1�xHgxSfilms was carried out with an emphasis on the difference between the �- and �-phase alloys deposited at 30°C as described earlier [27]. A linear increase inbandgap (up to 0.9 eV for x � 0.33) for the �-phase alloy and a linear decrease ofthe �-phase down to ca. 0.18 eV for x � 0.33 was measured. The resistivities ofthe alloys (x � 0.14) were higher than for the pure PbS (ca. 10 �-cm) by a factorof ca. 5 (�-phase) and of 10 (�-phase). Both the photoconductivity response andthe thermoelectric power of the alloys were greater than for the pure PbS. Theelectrical properties were believed to be controlled mainly by intergrain barriers.

Using the FeCl3-containing solution, epitaxial films of �-Pb1�xHgxS weregrown on (111) Si or Ge single crystals, where x varied between 0 and 0.33 [28].The conditions to obtain epitaxy were low temperature (�20°C), relatively dilutesolution (concentrations not given, but the typical concentrations were high—metal concentration probably several hundred mM), and relatively thin films(�80 nm; above this thickness, �-phase deposition occurred). The films on Gewere (111) oriented, while those on Si were (112). The requirement for low tem-perature and relatively low concentration of reactants, both of which slow the de-position process, suggest that the epitaxy occurs if enough time is allowed forcrystal growth to occur. Decreased temperature will decrease both crystal growthand the rate of attainment of epitaxy (probably a surface diffusion process); forepitaxy to be preferred at lower temperatures implies that the effect of tempera-ture on the former is greater than on the latter. The attainment of epitaxial growthis strong evidence for an ion-by-ion mechanism, even though parallel homoge-neous precipitation occurred in the solution.


Terminal thickness usually decreases with increase in temperature, due tofaster homogeneous precipitation in the solution in parallel with film deposition.In this study, the terminal thickness increased with increasing temperature. Thistrend also suggests that ion-by-ion growth dominates, since homogeneous precip-itation is less likely in an ion-by-ion process than in a cluster one. Although suchhomogeneous precipitation was observed, it seems likely that it occurred to alesser extent than it would in a cluster process.

8.2.7 (Cd,Bi)S

Films of BiMCdMS were deposited from triethanolamine/ammonia-complexednitrates of Cd and Bi using thiourea [29]. The films were deposited at a pH of ca.10 on glass or Si (111) at 95°C (for 90 min), followed by cooling to room tem-perature and continuing deposition for 24 hr. Rutherford Backscattering (RBS)analysis showed that both Cd and Bi were incorporated into the films very roughlyin proportions similar to those present in the solution. No structural data was givento support solid solution formation. The gradual variation of bandgap with com-position of annealed films (450°C in argon for 2hr), from 1.65 eV (for pure Bi2S3)to 2.43 eV (for pure CdS), suggested solid solution formation, at least for the an-nealed films.

8.2.8 (Cu,Pb)S

Two-phase films of PbSMCuS were deposited on glass from a triethanolamine/thiourea bath at room temperature [30]. As deposition proceeded, the films be-came Pb-rich as Cu was depleted by more rapid formation of sulphide. The resis-tivity of the films was �10 �-cm (1 m�/sq. for a film thickness of ca. 0.1 �m).

8.2.9 (Cu,Bi)S (and (Cu,Sb)S)

Films were deposited from a triethanolamine/thiourea bath containing CuSO4 andBi(NO3)3 [31]. No compositional or structural characterization was given; there-fore there is no evidence that this was a solid solution or even a mixture of Cu andBi sulphides.

An example of solid solution formation by separate deposition of binary lay-ers followed by annealing to interdiffuse the two layers is given for Cu3BiS3 de-position [32]. Bi2S3 (film thickness ca. 90 nm) was deposited at room temperaturefrom a Bi(NO3)3/triethanolamine/thioacetamide bath onto glass slides. CuS(300–600 nm thick) was then deposited on this film from a CuCl2/tri-ethanolamine/ammonia/NaOH/thiourea bath at room temperature. The films wereannealed at 250°C for 1 hr. Formation of the Cu3BiS3 phase could be seen fromthe XRD pattern. Measurement of precipitated powders (prepared by putting theBi2S3 precipitated in the first deposition in the CuS deposition solution) annealedat 300°C showed more clearly the formation of the solid solution.


Transmission spectra of the annealed films showed an approximate bandgapof 1.9 eV. The films were p-type semiconductors with fairly low resistance (de-pending on annealing conditions; even the as-deposited films had a sheet resis-tance of 7 k�/sq., which dropped to ca. 100 �/sq. on mild annealing at 150°C).

In a similar manner, Cu3SbS4 was formed by depositing a layer of CuS (asearlier) onto a previously deposited film of Sb2S3 (from a thiosulpahte bath) andannealing in N2 at 250°C [33]. These films (typically between 0.1 and 0.3 �mthick) were highly conducting (some tens of �/sq). The films were evaluated forsolar-control purposes and exhibited good IR reflectivity/low IR transmittancewith sufficient visible transmittance.

8.2.10 (Pb,Sn)Se

Films were deposited from solutions of lead and tin salts (the salts used were notspecified) with ammonium acetate, ethylenediamine, and selenourea at a pH � 9(probably at least 11) [34]. To obtain thicker films, deposition was repeated anumber of times and the films were annealed; therefore it is not known if solid so-lution formation occurred in as-deposited films. In annealed films, Pb1�xSnxSesolid solutions with x up to 0.11 were verified by XRD. The spectral response ofthe photoconductivity of the (annealed—as-deposited films were not photosensi-tive) films shifted from a peak at ca. 4 �m (pure PbSe) to ca. 7.5 �m (11% Sn),supporting solid solution formation of the annealed films. The room-temperature,dark resistance of the (probably annealed, but not certain) films varied from 1 to300 k�, depending on deposition conditions.

8.2.11 (Pb,Bi)S

From solutions of Bi and Pb nitrates, complexed with triethanolamine and ammo-nia, mixed sulphides were deposited with thiourea on glass at pH values between9.5 and 11 and at 100°C (initially) followed by slow cooling in the solution [35].Elemental analyses showed the presence of both metals in the films. It is not clearwhether solid solution formation occurred in the as-deposited films, although thelattice parameters did vary non-monotonically, depending on composition.

8.2.12 CuInS2 and CuInSe2

CuInS2 (and, even more, CuInSe2) are strong candidates for thin-film photovoltaiccells. For this purpose, the chalcopyrite structure (which is an ordered lattice) ispreferred over the disordered, zincblende form. Due to the large absorption coef-ficients of these materials, a 1-�m-thick film is more than enough to absorb al-most all the suprabandgap radiation. Somewhat thicker films are generally used,due to problems of pinholes, which commonly occur in thinner films. A numberof methods have been used to deposit these films. Surprisingly, very few (pub-lished) attempts have been made to deposit them by CD.


A note of caution is necessary when dealing with these materials. It is nottrivial to distinguish between CuInS(Se)2 and some phases of CuMS(Se). Diffrac-tion and optical properties may be similar. Elemental analysis is particularly im-portant to verify inclusion of indium in the films and in the correct ratio. A fin-gerprint of the chalcopyrite XRD is the presence of a weak peak at 2� � 17–18°,corresponding to the (101) chalcopyrite reflection. This is often not seen, althoughthis could be either because the deposit is not chalcopyrite or because weak peaksare usually not seen in nanocrystalline materials with particularly small crystalsize.

8.2.12.1 (Cu,In)S

Deposition of single-phase, chalcopyrite CuInS2 was claimed using a solution ofCuCl2, InCl3, complexed by triethanolamine and ammonia, with thiourea [36].(Note that this claim was contested, based on the diffraction and compositionaldata [37,38].) The best films were somewhat Cu rich (typically Cu1.08In1S1.5) and,more importantly, very S deficient. The excess Cu is not surprising, considering(1) the greater concentration of Cu in the solution compared to In and (2) the lowersolubility of CuMS compounds compared to InMS ones (see, however, the dis-cussion on CuMInMS compound precipitation in Sec. 8.1). The fact that filmswith excess In can be obtained with In:Cu ratios in solution of less than 1 suggeststhat the coprecipitation is more complicated than expected based solely on solu-bility products or even taking into consideration adsorption of In on CuMS as in-ferred from early studies on this and similar systems described in Section 8.1 [3,4].

Overall, it appears likely that the films contained chalcopyrite CuInS2

mixed with other phases with similar diffraction patterns. Separate microstruc-tural characterization (EDS) of films with varying composition (ca. 10% excessCu or In) showed the formation of separate phases of Cu2S and In2S3, respec-tively, along with the CuInS2 [39]. The best films were obtained at high deposi-tion temperatures (80°C) and with stirring. Lower deposition temperature resultedin poorer stoichiometry (less S), and stirring improved film uniformity. Grain size,measured by TEM (which does not necessarily show crystal size) was 100–400nm.

From the optical transmission spectra, a bandgap of 1.50 eV was found forthe most stoichiometric (in terms of Cu:In � 1; the S content was always found tobe low) films. This value dropped slightly for nonstoichiometric films [39]. Re-sistivities varied with composition, from ca. 50 �-cm for In-rich films down to ca.0.1 �-cm for very Cu-rich ones [36].

8.2.12.2 (Cu,In)Se

Three studies used essentially the same baths—ammonia and triethanolamine-complexed Cu and In salts—and selenosulphate as Se source [40–42]. In all cases,


a Cu:In ratio close to unity in the solution resulted in an optimum (and ca. unity)ratio in solution. In fact, the ratio in the film was not strongly dependent on the ra-tio in solution (at least, over the narrow range measured) [42]. This is not obviousbased on solubility product considerations and suggests some form of compoundformation, as described in the related precipitation experiments of Rudnev et al.[3,4].

In the earliest study, the deposition was carried out at room temperature; noelemental analysis was made, and the diffraction data do not show the presence ofchalcopyrite CuInSe2, although the sphalerite phase might be present. The otherstudies used higher temperatures, a parameter that appears to be important. Morestoichiometric films were found at 90°C than at 50°C [41], and in one case, good,adherent films were obtained only at 85°C; at lower temperatures, powdery, non-adherent films were obtained [42]. Both these studies reported chalcopyriteCuInSe2 as deposited. A bandgap of 1.08 eV was calculated from the optical spec-trum for the low-temperature deposition of Ref. 40 and one of 0.9 eV for the high-temperature (90°C) deposition of Ref. 41. All films that were characterized forconductivity type were found to be p-type, with resistivities that varied between0.08 and 500 �-cm, depending mainly on the excess Cu content, which resultedin low-resistivity films [41,42]. These are low values and suggest appreciable Cuexcess, although elemental analysis showed some of the films to be close to stoi-chiometric.

A similar deposition, using ammonia but with citrate complexant instead oftriethanolamine and at 40°C, was also reported [43]. From XRD measurements(not shown in the study), predominantly chalcopyrite CuInSe2 was reported if theCu:In:Se ratio in solution was ca. 1:1:2. EDS analysis confirmed this approximateratio in the films. From the absorption spectra, a bandgap of 1.4 eV was measured,which decreased to 1.15 eV on annealing to 520°C (the literature bandgap ofCuInSe2 is ca. 1.0 eV).

Resistivity and Hall measurements of these films as a function of composi-tion are interesting. The resistivity increased to a sharp maximum of ca. 108 �-cmat a Cu:In ratio of 1.5; lower values of Cu:In resulted in lower-resistance p-typefilms, while higher values (more Cu) gave low-resistance n-type films (Fig. 8.3).This is unexpected in that (a) In-rich films of CuInSe2 are normally n-type, whileCu-rich films are p-type and (b) the highest resistivity would be expected forCu:In � 1 if the material is CuInSe2. These results, together with the higher-than-usual bandgap of the (mainly annealed) films, suggest that the films are not sim-ply single-phase CuInSe2, but either a mixture of phases or a different composi-tion.

The most recent investigation is closest to the previous one, but using onlycitrate as a complexant (no ammonia), a lower pH than the other studies (8 insteadof ca. 10), and deposited at room temperature [44]. The films from this depositionwere not adherent as deposited and required annealing (300°C) to become adher-


ent. The films again appeared, from XRD, to be chalcopyrite. The XRD peaks ofthe as-deposited films were rather broad (crystal size of ca. 14 nm). As with theprevious study, the bandgap of the as-deposited films was anomalously high (1.3eV), which, considering the small crystal size, may be a quantum size effect. Af-ter annealing, the crystal size almost doubled and the bandgap dropped to the nor-mal CIS value of 1.02 eV. The films were p-type and highly conducting: 10�3 �-cm as deposited and ca. 2 �-cm after annealing. Again, these low resistivitiessuggest that free CuMSe species were present.

8.2.13 (Cd,Sn)O

Cadmium stannate (Cd2SnO4) was deposited from a solution of CdCl2 and SnCl4using NH4F, ostensibly as a freezing agent, although it is probable that it alsofunctioned as a complexant (see this technique for SnO2 deposition in Chap. 7)along with a small amount of AgNO3 as catalyst (not clear for what) and NaOHto adjust pH to between 7.5 and 8.5 [45]. The film grew to a maximum thicknessof 0.8 �m in 40 min at a pH of 7.5 (faster deposition but lower terminal thicknessat higher pH). X-ray diffraction of films annealed at 200°C or above showed themto be Cd2SnO4 with a grain size of 25 nm (20 nm before annealing).

FIG. 8.3 Resistivity vs. Cu:In ratio for CuMInMSe films (annealed at 520°C). (Adaptedfrom Ref. 43 with permission from Elsevier Science.)


The optical transmittance and reflectance spectra of the as-deposited filmand films annealed in H2 and in vacuum at ca. 200°C are shown in Figure 8.4. Allthe films show high transmission in the visible/near-IR region, in particular the an-nealed ones. A high transmittance is important if the films are to be used for theirtransparent, conducting properties. The blue shift of the absorption after anneal-ing was explained by the increase in bandgap due to a high free-carrier concen-tration (the Moss–Burstein shift caused by filling of the lower part of the conduc-tion band by free carriers). The increased carrier concentration after annealing isalso the cause of the shift in the reflectance spectrum to shorter wavelengths uponannealing; annealing in H2 results in a larger free-carrier concentration due to re-moval of more oxygen (heavier n-type doping). The bandgaps, calculated from thetransmission spectra, were all indirect and were 2.7 eV (as deposited), 3.1 eV(vacuum-annealed), and 3.2 eV (H2-annealed).

The resistivities of the films were 4 � 10�1 (as deposited), 10�2 (vacuum-annealed), and 4 � 10�3 �-cm (H2-annealed). The decrease in resistivity was duemainly to increase in free-electron concentration (2 � 1018, 5 � 1019, 1020 cm�3)for the three films; the mobility increased by a factor of two between the as-de-

FIG. 8.4 Transmittance and reflectance spectra of CdMSnMO (Cd2SnO4) films: as de-posited, annealed in H2 and in vacuum, both at ca. 200°C. (Adapted from Ref. 45.)


posited and H2-annealed films. The increase in electron concentration was due tooxygen removal, as mentioned earlier. The increased mobility was possibly due todesorption of surface-adsorbed oxygen.

8.3 MIXED CHALCOGEN COMPOUNDS

A number of mixed sulphide/hydroxides have been deposited, mainly in thesearch for improved window layers for photovoltaic cells (Chap. 9). These aremostly probably mixed-phase films, although in one case, In(OH)S, experimentalevidence suggests true compound formation [46]. Most of these films have beendealt with in previous chapters (see Chap. 4 under ZnS and Chap. 6 under In andSn sulphides). One study (described from the viewpoint of its properties in pho-tovoltaic cells in Chap. 9) has not been described previously and will be men-tioned briefly here. This deals with Zn(O,OH) and Zn(O,OH,S) deposited fromZn-ammine solutions, the latter film from solutions also containing thiourea [47].It is of interest to note that the Zn(O,OH) films did not deposit on glass but did onboth ZnO- and CuInSe2-type substrates. Even after annealing at 300°C, hydrox-ide groups were still present in those films.

8.3.1 Cd(S,Se)

In two of the studies made of Cd(S,Se) deposition [48,49], the solutions and con-ditions of deposition were similar. In both cases, Cd was complexed with ammo-nia, a mixture of thiourea and selenosulphate was used, and the deposition wascarried out on glass at 75°C.

From the study of Kainthla et al. [48], XRD of the films showed clearly thatsolid solution formation occurred; the (predominantly sphalerite) diffractionpeaks shifted with change in composition. For compositions with S concentration� 60%, only zincblende structure formed; the amount of wurtzite increased withincreasing S content but was always low. The concentration of S in the films wassomewhat greater than that in the deposition solution; i.e., S deposited preferen-tially. This is not surprising since CdS deposition is normally faster than that ofCdSe. The concentration of ammonia was increased as the thiourea:selenosul-phate ratio increased, ostensibly to slow down the rate of formation of CdSthrough decreased Cd2� concentration (although the rate of CdSe formation isalso dependent on this same factor).

Optical spectra showed a gradual shift of the onset with composition, as ex-pected for a true solid solution. Additionally, the refractive index (and thereforethe dielectric constant) increased gradually as the Se content increased, mirroringthe larger dielectric constant of CdSe compared with CdS.

The study of Ref. 49 gave similar results, although XRD peaks of the filmswere very weak and difficult to interpret. In addition, the electrical resistivity of


the films was measured. The resistance of the S-rich and slightly Se-rich films wastypically 1–3 � 106 �-cm, while for high Se concentrations the resistivity fell toa value of 7 � 104 �-cm for pure CdSe.

Another study of Cd(S,Se) deposition was similar to the preceding ones,with the differences that triethanolamine was used instead of ammonia (solutionpH � 10.4) and the deposition was carried out at 55°C [50]. The main differencein the films was that a true single-phase solid solution formed over only part of thecomposition range (for values of x in CdSexS1�x between 0 and 0.4 and between0.85 and 1). This was paralleled by sharp absorption onsets for the solid solutionsand more gradual ones for the mixed-phase systems. Crystal sizes varied in a non-monotonic manner from 27 nm (CdS) to 9 nm (CdSe). The resistivities were of thesame order (for CdS, slightly less) as those of the previous study, although thevariation of resistivity with composition was somewhat different. However, sincethe variation in resistivity over the entire composition range was only a little morethan an order of magnitude, which is not particularly large, such differences neednot be very meaningful.

The same group also studied In-doped (by addition of InCl3) CdS1�xSex

[51]. As the In concentration increased, the degree of crystallinity (measured bythe height of the XRD peaks) and crystal size increased, reached a peak, andthen decreased. These structural changes correlated with other properties: Thebandgap and resistivity were minimum when the crystallinity and crystal sizewere maximum. The concentration of In in the films was much higher than thatin solution (�2% of the Cd concentration, compared to 0.1% in the solution)[52].

8.3.2 Zn(S,Se)

Zn(S,Se) has been deposited on both glass and on single-crystal GaAs (110) froma hydroxide-complexed solution of Zn2� using, as for Cd(S,Se), a mixture ofthiourea and selenosulphate [53,54]. Apparently conditions were chosen to givethe composition ZnS0.056Se0.944 because of its perfect lattice match with the GaAssubstrate. The composition did not appear to be dependent on the deposition tem-perature.

Room-temperature deposition resulted in films with very broad peaks,which sharpened considerably with increasing deposition temperature to give acrystal size of ca. 20 nm at a deposition temperature of 90°C. The high-tempera-ture films on GaAs exhibited a fairly high degree of epitaxy, as seen by the spotsin the electron diffraction pattern.

8.3.3 Pb(S,Se)

Pb(S,Se) was deposited from a hydroxide-complexed solution of Pb(NO3)2 us-ing, as before, a mixture of thiourea and selenosulphate [55]. The films were


confirmed to be single phase by XRD (Debye–Scherrer photographs). In con-trast to the case for Cd and Zn, Se was preferentially deposited. (Small seleno-sulphate concentrations in solution resulted in much larger Se concentrations inthe film. For example, from a solution containing 0.5 mM selenosulphate and500 mM thiourea, the film composition was ca. PbS0.75Se0.25.) This was ex-plained by the large difference in solubility products of PbS and PbSe (nearly 10orders of magnitude; see Table 1.1). For Cd and Zn, this difference is some or-ders of magnitude less. Thus, while for Cd and Zn it seems that the faster de-composition of thiourea compared to selenosulphate more than compensates forthe lower-solubility products of the selenides, for Pb the difference in solubilityproducts between sulphide and selenide becomes the main composition-deter-mining factor.

From XRD line widths, it was noted that the crystal size increased as the Sconcentration increased, although values were not given. A simple but usefulcharacterization technique to quantify film thickness homogeneity was used here.The transmittance of a focused light spot scanned across the sample showed ex-cellent homogeneity.

8.3.4 Bi2(S,Se)3In this case, thioacetamide was used as the sulphur source, instead of thiourea asfor the previous mixed sulphides-selenides (selenosulphate, as before, was used asthe Se source) [56]. Bi(NO3)3 was complexed with triethanolamine and the pH ad-justed with ammonia to 8.2. The deposition was carried out at 55°C. The compo-sition was varied by varying the thioacetamide/selenosulphate ratio. Although itis not clear what the elemental compositions of the various films were, from thelimited XRD data given, it seems that solid solution did occur. The crystal sizesincreased from 6 nm (pure sulphide) to 13 nm (pure selenide), and bandgap val-ues decreased over the same range from ca. 1.9 to 1.0 eV.

8.4 BILAYERS OF DIFFERENT SEMICONDUCTORS

This chapter has dealt with true ternary compounds, with the underlying implica-tion that deposition of separate phases is undesirable. However, it needs to bestressed that what is undesirable for one purpose may be preferable for another(examples being small crystal size and a large amount of scattering). So, too, acomposite of different phases may be the goal of a particular deposition. This is-sue does not appear to have been dealt with in CD.

Taking the principle of separation of phases one step further, separate lay-ers may also be deposited, one on top of another. This has been done in a numberof cases and should present no problem (taking into consideration that there maybe some cases where deposition of the second layer will destroy or change in some


way the first one). There are, however, two examples in the literature of deposi-tion of a bilayer from a single deposition solution.

In the first [57], a bilayer of CdS/ZnS was formed using the electrochemi-cally assisted technique described for CdS (Sec. 4.1.6.6). In this technique, a mix-ture of CdCl2 and ZnCl2 was mixed with thioacetamide at a pH of 2.45 (addedHCl) and deposited on ITO/glass at 70°C. In the absence of electrochemical po-larization, only ZnS, with a low percentage of CdS, deposited after an inductionperiod. CdS does not deposit by itself under these acidic conditions. When sub-jected to cathodic polarization (�0.65 V vs. S.C.E.), CdS preferentially depositeddue to local increase in pH at the cathode (by water electrolysis), while eventually,as the Cd concentration dropped, ZnS formation became more favorable. Depthanalysis of the films showed that Zn was formed preferentially at the surface whileCdS formed preferentially near the substrate. This order is expected, consideringthe lower-solubility product of CdS.

The second example, which was a pure CD process, produced ZnO on topof CdS [58]. The principle is based on the facts that CdS deposits much more read-ily than does ZnS (see Sec. 4.4.1) and that ZnO (or Zn(OH)2, which readily con-verts to ZnO) tends to deposit readily, more so than ZnS unless under conditionsof high active sulphur concentration (whether sulphide ion or sulphur-containingcomplex) and low pH. The solution contained Cd and Zn ions (the latter in excess)complexed with ethanolamine and ammonia (therefore at least fairly high pH) andthiourea. Cross-sectional microprobe analysis showed that the film contained themore readily deposited CdS at the substrate (ca. 0.3 �m thick), covered with a 2-�m ZnO layer, which formed as the solution conditions (probably mainly the dropin Cd concentration) favored the ZnO deposition. The films were pale yellow, andoptical spectroscopy showed two transitions—one at the CdS bandgap (ca. 2.6 eV,greater than the bulk bandgap, suggesting that the CdS crystal size was small—ca.4–5 nm) and the other corresponding to ZnO at 3.2 eV. The films were photosen-sitive, about an order of magnitude more sensitive than ZnO deposited by itself.

A different form of bilayer can be formed using topotactic exchange reac-tions. This type of exchange is well known, e.g., for the conversion of CdS intoCu2S by immersing in a hot CuCl solution, used in the past for fabricatingCdS/Cu2S photovoltaic cells (see Sec. 9.1.2). It has been used more recently toconvert CD films of one semiconductor into another, e.g., CdS and CdSe intoAg2S and Ag2Se [59] and SnS2 into Ag2S [60]. While these studies describe con-version of one semiconductor into another, it is clear that, if carried out in a con-trolled fashion, partial exchange can occur, leading to the expected formation of ashell of the exchanged semiconductor around a core of the original semiconduc-tor for each individual crystal in the film (assuming the film to be at least some-what porous, as it invariably will be). This process therefore can lead to films ofcore/shell nanocrystals.


8.5 CONCLUSION

It is fair to state that the understanding of deposition of ternary compounds lagsbehind that of binaries. A better understanding of the factors that control codepo-sition, as well as solid solution formation, is needed. However, it is also clear thatthere is scope for deposition of a wide range of compounds, not only ternaries, butquaternaries and even higher-multinary materials. Additionally, the scope for de-position of mixed-phase films, either as consecutive layers (as shown earlier) oras composites, is great, and this aspect of CD will undoubtedly be pursued.

REFERENCES1. CS Brooks. Metal Finishing 88:21, 1990.2. IM Kolthoff, DR Moltzau. Chem. Rev. 17:293, 1935.3. NA Rudnev, GI Malofeyeva. Talanta 11:531, 1964.4. NA Rudnev, IV Melikhov, AM Tuzova. Zh. Anal. Khim. 28:635, 1973.5. KL Chopra, RC Kainthla, DK Pandya, AP Thakoor. In: Physics of Thin Films. Vol.

12, Academic Press, New York and London, 1982, p 167.6. G Contreras-Puente, O Vigil, M Ortega-Lopez, A Morales-Acevedo, J Vidal, ML Al-

bor-Aguilera. Thin Solid Films 361:378, 2000.7. JM Doña, J Herrero. Thin Solid Films 268:5, 1995.8. NC Sharma, RC Kainthla, DK Pandya, KL Chopra. Thin Solid Films 60:55, 1979.9. T Nakazawa, S Kuranouchi, A Ashida, N Yamamoto. In: 12th EC Photovolt. Solar

Energy Conf. Amsterdam, the Netherlands, 1994, p 601.10. GK Padam, GL Malhotra, SUM Rao. J. Appl. Phys. 63:770, 1988.11. SA Al Kuhaimi, Z Tulbah. J. Electrochem. Soc. 147:214, 2000.12. T Yamaguchi, Y Yamamoto, T Tanaka, Y Demizu, A Yoshida. Thin Solid Films

281–282:375, 1996.13. T Yamaguchi, Y Yamamoto, T Tanaka, A Yoshida. Thin Solid Films 344:516, 1999.14. KC Sharma, JC Garg. J. Phys. D: Appl. Phys. 23:1411, 1990.15. DS Sutrave, GS Shahane, VB Patil, LP Deshmukh. Mater. Chem. Phys. 65:298, 2000.16. ES Rittner, JH Schulman. J. Phys. Chem. 47:537, 1943.17. M Skyllas-Kazacos, JF McCann, R Arruzza. Appl. Surf. Sci. 22/23:1091, 1985.18. LP Deshmukh, KM Garadkar, DS Sutrave. Mater. Chem. Phys. 55:30, 1998.19. BB Nayak, HN Acharya, GB Mitra. Bull. Mater. Sci. 3:317, 1981.20. BB Nayak, HN Acharya. J. Mater. Sci. Lett. 4:651, 1985.21. GB Reddy, DK Pandya, KL Chopra. Sol. Energy Mater. 15:383, 1987.22. GA Kitaev, VF Markov, LN Maskaeva, LE Vasyunina, IV Shilova. Inorg. Mater.

26:202, 1990.23. GA Kitaev, LN Maskaeva, VF Markov, AY Kurkin, LE Vasyunina. Inorg. Mater.

25:1065, 1989.24. LN Maskaeva, VF Markov, GA Kitaev. Russ. J. Appl. Chem. 73:751, 2000.25. NC Sharma, DK Pandya, HK Sehgal, KL Chopra. Mat. Res. Bull. 11:1109, 1976.26. NC Sharma, DK Pandya, HK Sehgal, KL Chopra. Thin Solid Films 42:383, 1977.


27. NC Sharma, DK Pandya, HK Sehgal, KL Chopra. Thin Solid Films 62:97, 1979.28. NC Sharma, DK Pandya, HK Sehgal, KL Chopra. Thin Solid Films 59:157, 1979.29. S Misra, HC Padhi. J. Appl. Phys. 75:4576, 1994.30. R Suarez, PK Nair. J. Solid State Chem. 123:296, 1996.31. LP Deshmukh, DS Sutrave, BM More, CB Rotti, KM Garadkar. Semicond. Devices

421, 1996.32. PK Nair, L Huang, MTS Nair, H Hu, EA Meyers, RA Zingaro. J. Mater. Res. 12:651,

1997.33. MTS Nair, Y Pena, J Campos, VM García, PK Nair. J. Electrochem. Soc. 145:2113,

1998.34. VM Markov, LN Maskaeva, LD Loshkareva, SN Uimin, GA Kitaev. Inorg. Mater.

33:555, 1997.35. N Parhi, BB Nayak, BS Acharya. Thin Solid Films 254:47, 1995.36. GK Padam, SUM Rao. Sol. Energy Mater. 13:297, 1986.37. D Cahen. Sol. Energy Mater. 15:225, 1987.38. GK Padam, SUM Rao. Sol. Energy Mater. 15:227, 1987.39. GK Padam, GL Malhotra, SUM Rao. Phys. Status Solidi (a): 109:K45, 1988.40. RN Bhattacharya. J. Electrochem. Soc. 130:2040, 1983.41. GK Padam. Mat. Res. Bull. 22:789, 1987.42. KR Murali. Thin Solid Films 167:L19, 1988.43. JC Garg, RP Sharma, KC Sharma. Thin Solid Films 164:269, 1988.44. PK Vidyadharan Pillai, KP Vijayakumar, PS Mukherjee. J. Mater. Sci. Lett. 13:1725,

1994.45. D Raviendra, JK Sharma. J. Appl. Phys. 58:838, 1985.46. D Hariskos, M Ruckh, U Ruhle, T Walter, HW Schock, J Hedstrom, L Stolt. Sol. En-

ergy Mater. Sol. Cells 41–2:345, 1996.47. K Kushiya, T Nii, I Sugiyama, Y Sato, Y Inamori, H Takeshita. Jpn. J. Appl. Phys.

35:4383, 1996.48. RC Kainthla, DK Pandya, KL Chopra. J. Electrochem. Soc. 129:99, 1982.49. RS Mane, CD Lokhande. Thin Solid Films 304:56, 1997.50. GS Shahane, BM More, CB Rotti, LP Deshmukh. Mater. Chem. Phys. 47:263, 1997.51. GS Shahane, KM Garadkar, LP Deshmukh. Mater. Chem. Phys. 51:246, 1997.52. GS Shahane, LP Deshmukh. Mater. Chem. Phys. 70:112, 2001.53. GN Chaudhari, S Manorama, VJ Rao. Thin Solid Films 208:243, 1992.54. GN Chaudhari, SN Sardesai, SD Sathaye, VJ Rao. J. Mater Sci. 27:4647, 1992.55. YS Sarma, HN Acharya, NK Misra. Thin Solid Films 90:L43, 1982.56. AR Patil, VN Patil, PN Bhosale, LP Deshmukh. Mater. Chem. Phys. 65:266, 2000.57. K Yamaguchi, T Yoshida, T Sugiura, H Minoura. J. Mater. Res. 13:917, 1998.58. SJ Castillo, M Sotelo-Lerma, RA Zingaro, R Ramirez-Bon, FJ Espinoza-Beltran, R

Guillemette, MA Dominguez. J. Phys. Chem. Solids 62:1069, 2001.59. CD Lokhande, KM Gadave. Mater. Chem. Phys. 36:119, 1993.60. CD Lokhande, VV Bhad, SS Dhumure. J. Phys. D: Appl. Phys. 25:315, 1992.


9Photovoltaic andPhotoelectrochemicalProperties

A large number of studies on CD have been driven by two related potential uses:photoelectrochemical (PEC) cells, mostly the earlier studies, and, more recently,photovoltaic (PV) cells. This chapter is devoted to these two topics where CDfilms have been used.

9.1 PHOTOVOLTAIC CELLS

9.1.1 Introduction

Since almost all thin-film (CdTe- and CuInSe2-type) cells today use CD films(mostly CdS), there is no attempt here to be comprehensive regarding the litera-ture. Rather, studies that emphasize the CD film itself are discussed. Before dis-cussing the role of the CD layer specifically, a brief overview of the relevant cellswill be given.

There are three main thin-film PV cells under development at present: amor-phous Si, CdTe/CdS, and CI(G)S/CdS [CI(G)S refers to copper indium (gallium)selenide]. Of these, the last two are polycrystalline (as opposed to amorphous),and both normally employ CD CdS. Crystalline Si cells are not thin films, beingat least tens and usually hundreds of microns in thickness, compared to the fewmicrons of active thickness of the thin-film cells.

Schematic diagrams of the CdTe and CI(G)S cells are shown in Figure 9.1.The main difference in their construction is that the CdTe cell is a superstrate


(backwall) cell (illuminated through the conducting glass substrate), while theCI(G)S cell is a substrate (frontwall) cell (illuminated through the front surface).The CI(G)S cell is a development of the original CuInSe2/CdS (CIS/CdS) cell,with Ga added to increase the bandgap. Pure CuGaSe2/CdS has also been investi-gated, although considerably less than CI(G)S, due to its (at present) considerablylower conversion efficiency. In the following, CIS will be used where CuInSe2 isintended, while CI(G)S refers to CuInxGa1�xSe2, with 1 � x typically 0.2 � 0.1.Some studies have also been made on CuInS2, which has a higher bandgap thanCIS and in principle should give a better cell (in practice it is inferior, althoughsomewhat better than CGS).

In both cells, the absorber layer (CdTe or CI(G)S) is a few microns thick,while the CD CdS (or other CD layer) is typically 50–100 nm thick. The CD layeris often called the “buffer” layer, a term that serves to show the lack of under-standing of its role. Nominally, the CdS is the n-type part of the p-n junction. Thebasic mode of action of a p-n PV cell is shown in Figure 9.2. The short-circuit cur-rent (denoted in this chapter as ISC) is the current flowing in the illuminated cellwhen the two sides (terminals) of the cell are shorted (full band bending, left fig-ure) while the open-circuit voltage (denoted as VOC) is the voltage generated be-tween the terminals at open circuit when no external current flows (right figure).The absorber CdTe and CI(G)S are always p-type in these cells; n-type absorbershave been little investigated, mainly because suitable high-bandgap p-type mate-rials are not readily available. A high bandgap of the buffer layer is necessary be-cause light passes through this layer on the way to the absorber, and some of thelight absorbed in the buffer layer is lost for current generation. The superiority ofCD CdS over evaporated CdS in both types of cell suggests that something otherthan loss of photons absorbed in the buffer layer is involved. The factors believedto contribute to this superiority, in particular the effects of the deposition solutionon the absorber, are discussed later.

FIG. 9.1 Schematic cross section of CI(G)S/CdS and CdTe/CdS PV cells.


The foregoing explanation of the operation of these cells, while very basic,will be almost sufficient for our purposes (there are many sources explaining themode of action of PV cells in more detail). One other process, which plays an im-portant role in PV cells in general, should be described: electron–hole recombina-tion. This is central to PV cell operation. Photogenerated electrons and holes areideally separated at the p-n junction and flow in opposite directions to give an ex-ternal current. However, there are many pitfalls awaiting these charges on theirway to the terminals where at least one of them can be extracted. These pitfalls,which cause the electrons and holes to recombine before external current flows,are various forms of recombination centers. They can occur at the interface be-tween the p- and n-type semiconductors, at grain boundaries, or in the bulk of thesemiconductor crystals. A major part of PV cell research is devoted to minimiz-ing such recombination centers.

9.1.2 CuxS/CdS Cells

Before describing studies on the CIS and CdTe cells, there are two CD-related pa-pers on the Cu2S/CdS cell, which was intensively investigated around 20 yearsago and was eventually abandoned because of perceived insoluble stability issues,a perception that, it should be noted, while widely held, is not undisputed. Shouldthis cell make a comeback, CD is likely to be a method that will be considered foreither of the two semiconductors or even for both.

One study utilized CD CdS, built up from several layers (probably a totalthickness of the order of a micron) and annealed [1]. The CuxS layer was formedby the usual (for this type of cell) topotactic reaction between a CuCl solution andthe CdS substrate. While the cell performance varied over a wide range, depend-

FIG. 9.2 Band diagrams of p-n cell in the dark (or under illumination at short circuit)and under illumination at open circuit.


ing on the CD process, the maximum efficiency obtained was 0.13%; all parame-ters of the cell were very poor.

In the other study, CuxS was chemically deposited on (presumably evapo-rated) CdS films from a triethanolamine/ammonia/thiourea bath (see Chap. 6,copper sulphides) [2]. Very low currents and poor fill factor were obtained, al-though the VOC was reasonable (ca. 0.5 V), with an efficiency of ca. 0.5%. The sto-ichiometry of the Cu-S was not ideal for PV cell use, although this could be var-ied to an extent by electrochemical treatment.

In view of the sensitivity of the Cu2S/CdS cell to the nature and phase of theCu-S, it is likely that much better performance can be obtained if an effort is madeto do so.

9.1.3 CdTe/CdS Cells

While CD CdS is commonly used for CdTe-absorber cells, there is relatively lit-tle work that emphasizes the CdS film.

A 13.4% efficient cell fabricated by close space sublimation of CdTe on CDCdS was reported in 1991 [3], followed by a 14.5% cell a year later by the samegroup [4]. The CdS thickness was between 50 and 150 nm. The cells were illumi-nated through the tin oxide/glass, which was used as the substrate for the CdS de-position, and this geometry has been used ever since for these cells.

The most comprehensive study of the effect of the CdS deposition parame-ters on the resulting CdTe/CdS appears to have been made for electrodepositedCdTe [5]. The most simple variable is CdS film thickness. Clearly a minimumthickness is required for junction formation and to prevent shunting from theCdTe through the CdS to the substrate (usually conducting glass). On the otherhand, an increase in this thickness leads to a decrease in ISC due to light absorp-tion in the CdS, which is clearly seen as a decrease in the short-wavelength re-sponse of the cell. The optimum CdS thickness was found to be ca. 70 nm, al-though good cells were also made with more than twice this thickness.

Other deposition parameters affected mainly VOC and fill factor rather thanISC. These included an increase in thiourea concentration and the use of buffered(ammonium ion, lower pH) solutions; both these factors resulted in higher S:Cdratios, therefore more stoichiometric films (CD CdS films are often Cd-rich; thisdoes not necessarily mean n-type doped but is more likely due to the presence ofother Cd species, e.g., Cd(OH)2). The use of chloride as an anion in solution ratherthan sulphate also gave better cells. It was believed that all these factors influencedthe nature of both the CdS and the CdTe films after annealing the cell. Specifi-cally, it was thought that small grain size and high defect density (the CdS was be-lieved to be polytype with a large density of stacking faults [6]) in the cubic CdSfilm was beneficial for the resulting recrystallization process and for intermixingbetween the CdS and CdTe during the recrystallization and phase change (to


hexagonal phase) of the CdS during the annealing step. In this respect, it was notedthat if the CdS was recrystallized prior to CdTe deposition (and the cell then re-annealed as usual), the resulting cells were very inferior in all output parametersto normal cells.

The VOC of electrodeposited CdTe on CD CdS cells was studied as a func-tion of the CdS deposition parameters [7]. While there were a number of differentvariables involved, it was clear that conditions leading to thicker films (ca. 180nm), such as lower pH, high thiourea:Cd ratio (or possibly higher thiourea con-centration), and repeated deposition of the films, resulted in the highest voltage(ca. 0.7 V). The short-wavelength response was poorer for thicker films, not onlydue to absorption in the CdS, but also at wavelengths longer than the CdS ab-sorption onset, suggesting recombination at interface states.

A clear effect of CD CdS on heterojunction formation has been shown forCdTe that was electrodeposited onto CdS films on single crystal (1̄1̄1̄) InP [8].CdTe electrodeposited directly onto the (1̄1̄1̄) InP shows some degree of epitaxybut also considerable polycrystallinity (the latter not surprising, considering the9.5% lattice mismatch between the two materials). If, however, a thin (20–30 nm)CdS film was chemically deposited onto the same InP surface, the epitaxy of CdTeelectrodeposited onto the CdS/InP was found to be very good. The CdS waslargely epitaxial with the InP (the lattice parameters of CdS and InP are veryclose), with ca. 15% polycrystallinity. Interestingly, the CdTe deposited on theCdS exhibited an even higher degree of epitaxy than that of the CdS itself, show-ing that the small but appreciable amount of polycrystalline CdS did not substan-tially degrade the epitaxy of the CdTe. It was suggested that the improvement inepitaxy due to the CdS was caused by a graded interface. Similarly, an XRD com-parison of CdTe electrodeposited onto SnO2/glass (activated by a cathodic treat-ment) or onto CdS chemically deposited on SnO2/glass showed better crystallinity(narrower XRD peak) for the CdTe deposited on the CdS and also better texture(only the (111) reflection was seen, while a small additional (220) peak was evi-dent for the CdTe deposited directly on the SnO2) [9].

9.1.4 CdS/CI(G)S Cells

9.1.4.1 General Considerations

A considerably greater body of work with more emphasis on the CD buffer layerexists for this cell. Much of this involves the specific effects of the CD process atthe interface, and this will be discussed in a later section.

An experimental measurement of the band lineup between the CdS and so-lar-grade polycrystalline CIGS has been made using contact potential difference(Kelvin probe) measurements in air [10]. This lineup is shown in Figure 9.3. Inparticular, it shows that no spike was found in the conduction band. The presenceof such a spike (believed to occur from previous studies either on single crystals


or on polycrystalline films in vacuum) would constitute a barrier to electron flowover the interface.

From the practical point of view, reuse of the deposition solution after fil-tering out the precipitated CdS and addition of fresh reagent was shown to haveno effect on the device properties of CdS/CIGS cells (see Sec. 4.1.6.10 for moredetails of the deposition) [11].

9.1.4.2 Chemical Reaction and Diffusion at theCdS/CI(G)S Interface

There are a number of studies on the effects of the CD process on the surface ofthe CIS or CIGS. In many of these studies, the absorber surface was treated withpartial CD solutions, in particular, ammonia or ammonia � Cd2� [12–19]. Thereare several reactions that occur during these treatments, and these will be dis-cussed in general before specific results from the different studies are treated.

Aqueous ammonia removes surface oxides from the CI(G)S, in particular,indium (possibly due to the tendency of the CIS surface to be enriched in In-O).In this respect, it can be considered as an etchant, although the etching is limitedto the near surface region and does not continue (although in one case, it wasfound that, if oxygen is present, etching can continue, presumably due to contin-ual oxidation of the freshly exposed surface [13]).

FIG. 9.3 Band offset between CIGS and CdS. (After Ref. 10.)


Cd2� dissolved in ammonia has a particularly strong effect on the absorbersurface. The effect of ammonia � Cd with respect to the In and Cu concentrations(decrease in Cu) is opposite to that of ammonia by itself (decrease of In), and itwas noted that the surface reaction with all components present in the CD bathwould depend on the relative kinetics of the various partial reactions. Probablymore important, however, is that Cd rapidly substitutes for Cu for between severalnanometers and 20 nm into the absorber. This has been found to occur on singlecrystals of CIS [15,16] as well as on polycrystalline films, and therefore is not sim-ply grain boundary diffusion, as might at first be suspected. In contrast to these re-sults, one study found no Cd substitution for Cu in CIS single crystals and relatedthis to the supposed absence of a Cu-deficient layer in the single crystals, com-pared to films [19]. It is not clear whether Cd indiffusion promotes Cu outdiffu-sion or vice versa. It should be noted that Cu is readily complexed by the ammo-nia at the surface and therefore is easily removed. Also, the ionic radii of Cd2� andCu� are almost identical, facilitating exchange. It is believed that the indiffusedCd type converts CIS to n-type and that the junction is a buried n-p one rather thanlocated at the CdS/CIS interface. There is also the question of how abrupt the junc-tion is; the Cu-poor surface region has been considered a separate phase—CuIn3Se5—known as an ordered vacancy compound (OVC). Exchange of Cu andCd was shown to be easier for single-crystal CuIn3Se5 than for CIS [16].

The effect of this Cd/NH3 treatment on the PV properties are very marked.While cells fabricated without a buffer layer [ZnO sputtered directly on theCI(G)S] are very poor, with all parameters very low, the same cells, but subjectedto the Cd/NH3 treatment before ZnO deposition, are very much better, and in factthe efficiencies are only a little lower than CD CdS cells, due to lower VOC (ISC isactually often higher due to the better blue response in the absence of CdS). Thisis a particularly important result since it shows that the main role of the bufferlayer is not related to the specific properties of the CdS itself, but rather to near-surface modification of the CI(G)S. Substitution of Zn for Cd in the Cd/NH3 treat-ment gave comparable results [15]. This is in contrast to the use of CD ZnS, whichwas inferior to that of CdS, although not necessarily by much (see Section9.1.4.5).

However, the presence of the CD CdS is still required in order to obtain op-timal efficiency, and therefore the CdS itself does play some role, possibly to pre-vent sputter damage to the absorber (although Cd/NH3-treated absorbers, whichthen have evaporated CdS deposited, are still not as good as CD CdS cells). Theconformal coverage of the irregular absorber surface by the CD CdS is anotherfactor often invoked. Incomplete coverage (e.g., pinholes) could lead to someshunting between the ZnO and CI(G)S, either due to the lack of CdS at the pin-holes and/or because of sputter damage at the pinholes. Also in connection withthe nature of the CdS itself, studies on the effects of CdS thickness and impuritycontent on CdS/CIGS cell parameters have been made [20]. The impurity content


was controlled by the concentration of thiourea in the bath (see Sect. 4.1.7 for de-tails) while maintaining the CdS thickness constant. After a very thin CdS filmhad been deposited, the main effect of increasing both thickness and impurity con-centration was to increase the VOC. This suggests that these effects are related tothe “bulk” properties of the CdS rather than to effects of the different solutions onthe absorber surface. It should be kept in mind, however, that the predeposited ul-trathin layer either may not have totally covered the CIGS or may have beenporous enough to allow contact between the second deposition solution and theCIGS. In connection with impurity considerations, a general property of CD CdSis its relatively large oxygen content. While it is not clear whether this is of anyimportance for its use in PV cells, it is often studied in this respect. The subject ofoxygen impurities in CdS is treated in some detail in Section 4.1.7.

Photoluminescence (PL) measurements to monitor the changes due to theCd/NH3 treatment of CIGS have been carried out by two groups. Both show ma-jor, though very different, effects of the treatment on the PL spectra. In one case[14], two shallow subbandgap peaks, attributed to donor–acceptor pair recombi-nation, were found in the nontreated CIGS. After treatment, the higher-energy peakwas quenched and a strong new, lower-energy peak appeared. This could be inter-preted as the removal of states that caused interface recombination, although thenature of the strong lower-energy luminescence was not understood. Notably, sim-ilar behavior was obtained when CdS was deposited, providing further evidencefor the dominating role of the Cd/NH3 treatment in the CdS buffer layer formation.

In the second study, the bandgap luminescence was found to increase 15times after the Cd/NH3 treatment [16]. This could be interpreted as a passivationof recombination centers. The luminescence from CD CdS/CIGS was stronger (byca. two times) than from CdS evaporated on Cd/NH3-treated CIGS and muchstronger (nearly 20 times) than from CdS evaporated on nontreated CIGS. Fromthese results, it was believed that both the Cd/NH3 treatment and the CdS deposi-tion were important, although, again, it appears that the major effect is from theCd/NH3 treatment.

On an In-rich CIS surface (one from which good-quality cells could bemade), the change in surface composition was followed upon deposition of CdSfrom a complete bath [12]. An important observation was that initially, while Cdwas found at the surface, the S concentration was low, and stoichiometric CdSonly formed later in the deposition. It was stressed, however, that since the depo-sition was begun at room temperature and then gradually heated to 60°C, it mightbe that the Cd exchange process occurred preferentially at low temperature, whereCdS deposition was still very slow, while the thiourea decomposition, which wasstrongly temperature dependent, occurred more readily as the deposition temper-ature increased. The initial Cd species were believed to act as nucleation centersfor CdS formation. This heating regime also slowed down precipitation in solu-tion and resulted in a lower terminal thickness (ca. 75 nm compared to �100 nm


for a preheated solution [13]), which would make the deposition time to obtain therequired thickness (ca. 50 nm) less critical.

An investigation of CdS deposition on CIGS (In � Ga rich—good PV cellquality) substrates was carried out using two different baths: a “standard” bath anda low-thiourea (3 mM) one [17]. Cd was preferentially incorporated into (onto) thesubstrate compared to S in the initial stages of the deposition, in agreement withthe previous study. In contrast to the previous study was the conclusion that nomajor preferential removal of any of the substrate constituents occurred (In, Ga,and Se native oxides were removed). In addition, no clear evidence of compoundformation between the Cd and substrate was found, although it is clear that Cd wasincorporated into the substrate. Although the low-thiourea bath resulted in CdSwith lower impurity (mainly N compounds) levels and more Cd incorporation intothe CIGS, no improvement in cell performance resulted from such baths. Whilethe nature of the Cd incorporation in the early stages of the deposition could notbe unambiguously defined, it was suggested that this Cd was in the form ofCd(OH)2, which converted to CdS, either as the thiourea slowly hydrolyzed to sul-phide or as decomposition of a Cd(OH)2–thiourea complex occurred, both ofwhich are very temperature dependent.

An explanation for the beneficial effect of the CD process on CI(G)S cellswas suggested based on the known effects of oxygen treatment on these cells [21].Annealing in oxygen removes Se vacancies, which in turn decreases recombina-tion at grain boundaries, surfaces and at the CdS/CI(G)S interface. In oppositionto this beneficial effect, the oxygen treatment also has been shown to reduce pos-itive charge, and therefore band bending at the interface, and to increase Cu dif-fusion into the bulk of the CI(G)S, thus reducing the acceptor doping; both theseeffects are detrimental to cell performance. The CD deposition was believed to re-store the positive charge to the interface (with which the deposition solution wasin contact) but not to the grain boundaries (where it did not reach) by creation ofCd on Cu vacancies (CdCu) and possibly also removal of oxygen on Se vacancies(OSe). This removes the detrimental effect of the oxygen treatment at the interfacebut not the beneficial effect at the grain boundaries.

9.1.4.3 Epitaxy of CdS on CIS

The surface cleaning of the CIS also affected the mode of deposition of the CdS.The CdS was found to grow to a greater or lesser extent of epitaxy on single-crys-tal (heteroepitaxial layer) CIS [22]. Very good epitaxy of cubic CdS was found forcyanide-treated CIS; somewhat lower epitaxy was found for ammonia-treated sur-faces and poorer epitaxy obtained for untreated surfaces that contained consider-able oxides. Additionally, the epitaxy was only obtained at higher deposition tem-peratures (�70°C); at lower temperatures, the growth was polycrystalline.

Epitaxial growth of cubic CdS {111} on the CIGS {112} was also found tooccur using a room-temperature bath, with gradual heating to 80°C 18,19]. Again,


the epitaxy was explained by ammonia cleaning of the CIGS followed by ion-by-ion growth.

9.1.4.4 CuGaSe2-based cells

CuGaSe2 (CGS) has also been studied as a PV material, although efficiencies ofcells based on this semiconductor are, at present, much lower than those usingCIS or CIGS. There is one study on the specific interaction between CD CdSand CGS [23]. Two different CD baths were used to deposit CdS in this inves-tigation: one at 60°C and the other at 80°C (the latter with a higher ammoniaconcentration to slow down the deposition). Several pronounced differenceswere found between the two baths, in spite of the relatively small difference be-tween them. For Cu-rich CGS, Cu-S inclusions in the CdS were formed in thehigh temperature bath, due to interaction between Cu in Cu-rich CGS and CdS,but not in the low-temperature bath. Such inclusions could lead to shunting. ForGa,In-rich films (which gave better cells), the 80°C deposition resulted in Se be-ing found in the CdS layer. Additionally, the higher-temperature CdS was lessdefected and formed a less defected interface with the CGS. The 80°C deposi-tion gave better PV properties for Ga-rich films (up to 9.3% efficiency). The60°C deposition, however, was better for the (poorer cell quality) Cu-rich films,which could be explained by shunting through the Cu-S inclusions in the 80°CCu-rich CGS devices. This investigation clearly shows the necessity for opti-mization of the CD process, not only for every specific absorber material, buteven for different types of the same absorber.

9.1.4.5 Cd-Free Buffer Layers

While most of the reported studies on CD buffer layers deal with CdS, there havebeen a number of attempts to chemically deposit other materials. There are sev-eral reasons for this. One is the desire, particularly prominent in Europe, to find aCd-free cell, for obvious environmental reasons (this clearly is relevant for theCIS and not for the CdTe cell). Another reason is that part of the light absorbed inthe CdS is lost for current generation (a point that does not seem to have been rig-orously investigated but that probably is due to a high recombination rate for holesphotogenerated in the CdS). A higher-bandgap material will therefore allow moreshort-wavelength photons to reach the CIS absorber and generate photocurrent.Since the CD process is believed to be largely responsible for the beneficial prop-erties of the CdS films, it has been anticipated that other CD layers would behavelikewise. Other possible factors are the band lineup between CIS and CD layer and(photo)conductivity (or maybe better, resistivity, of the buffer layer, since it ap-pears that a high resistivity is beneficial).

ZnS. ZnS, with a high bandgap of 3.7 eV, is an obvious choice for alter-nate buffer layers. ZnS films, prepared using thioacetamide from an ammonia-free


bath (see Sec. 4.4.2 for a description of the films) were deposited onto CIS films[24]. The higher short-circuit current expected for the higher-bandgap ZnS wasobtained, although the overall efficiencies were lower than for CdS deposited onsimilar substrates, due to lower photovoltages. Efficiencies as high as 9% wereobtained.

Most ZnS baths probably give films containing some hydroxide (at highertemperature, maybe oxide), as discussed in Section 4.4.1. The previous bath maybe an exception, since it was presumably carried out at relatively low pH. UsingZn(O,OH,S), efficiencies as high as 12.8% were obtained on CIGS substrates[25]. This efficiency was obtained after ca, 60-mn illumination and reversibly de-creased when kept in the dark. Spectral response measurements showed the ex-pected increase in short-wavelength response. Such an efficiency can be com-pared with the best reported value over 16% obtained at that time; however, sincethe highest reported values are considerably higher than those more routinely ob-tained, then the value of 12.8%, while lower than that obtainable using CdS, is nottoo far below it. Ideally, such experiments should be compared with state-of-the-art CdS deposited on (as closely as possible) identical substrates.

A somewhat more recent study reported 15.1% for ZnS/CIGS (comparewith 17.0% state-of-the art CdS on the same CIGS substrates) [26]. This studyshowed the differences in optimization of the cell, depending on whether CdS orZnS was used. Thus, while a 40 nm-thick ZnS layer led to a large improvement inall cell parameters, as was also the case for CdS, in contrast to CdS, increasingthickness of the ZnS to 90 nm reduced the cell performance, giving in particular avery low fill factor. This was explained by the higher resistivity of the ZnS (by twoorders of magnitude), compared to CdS. In addition, while a high-resistivity sput-tered ZnO layer was deposited on the CdS prior to the conducting ZnO:A1 to ob-tain maximum VOC, such a layer degraded the ZnS-buffer cell. Since this ZnO wassputtered in an oxygen-containing atmosphere, and no such degradation occurredfor the conducting ZnO sputtered in an oxygen-free environment, it was believedthat the degradation was caused by negative oxygen ions or energetic neutral par-ticles. Finally, while the short-wavelength response of the ZnS cell was better thanthat of the CdS cell, this improvement did not compensate for the lower quantumefficiency at longer wavelengths for the ZnS cell.

ZnSe. ZnSe, with a bandgap of ca. 2.7 eV, is another obvious substitute forCdSe. As with ZnS, there is a tendency for CD ZnSe to contain hydroxide.Zn(Se,OH) deposited from a selenourea bath was deposited on CIGS [27] (seeSec. 4.5.2 for details of the Zn-Se films). Efficiencies up to 13.7% (highest litera-ture value �17% using CdS) were found. Spectral response measurementsshowed the expected improvement in short-wavelength response. The optimalthickness of the CD layer was 7–8 nm; a layer ca. twice that thickness resulted ina drop in efficiency (to 10.4%), mainly due to a drop in fill factor, probably be-


cause of the higher resistivity of the thicker film. From XPS measurements, theband alignment could be estimated; this is shown in Figure 9.4. Tunneling of elec-trons through the large barrier at the junction between the CIGS and Zn(Se,OH)conduction bands was presumed to occur readily for the thin films but much lessso for the thicker ones.

ZnSe, deposited by the same method, was also used as a buffer layer forCuInS2 cells [28]. Higher currents and voltages but lower fill factor were obtained,compared to CdS, with a slightly lower overall efficiency. The band lineup for theZnSe/CuInS2 junction was also measured by XPS for this system.

ZnO. ZnO, which is normally used as the conducting window layer on theCIS-type cell, has also been used as a buffer layer for CuInS2 cells by annealingCD Zn(OH)2 (deposited from Zn2�/ammonia solutions) [29] (see Sec. 7.2.18,ZnO). The efficiency was much lower (3.8%) compared with comparable CdSbuffer layers (8.6%). The difference was due to much lower open-circuit voltageand resulting lower fill factor; the photocurrents were similar using both bufferlayers. The increase in photocurrent at short (�500 nm) wavelengths due to thehigher-bandgap ZnO was offset by lower photocurrents over the rest of the active

FIG. 9.4 Band diagram of the CIGS/Zn(Se, OH)/ZnO cell. (After Ref. 27.)


spectrum. While the thickness of the ZnO was not given, it was noted that the InXPS signal was still visible after the ZnO deposition. Comparing this with the pre-vious study by the same group using Zn(Se,OH), where the In peak disappearedwhen the CD layer was 7 nm and was already weak at ca. 4 nm, suggests eitherthat the ZnO layer was thinner than optimal or that the coverage was not homo-geneous. In view of the previous experiments on CI(G)S, which showed that aZn2�/ammonia treatment was as effective as a Cd2/NH3 treatment and both notmuch less effective than optimally deposited CdS, the Zn2�/NH3 bath used herewas considerably inferior to CdS buffer layers, even assuming an overly thin filmdeposited. Whether this difference is due to the different substrate (CuInS2 insteadof CI(G)S), to the annealing treatment, or to some other reason remains to be in-vestigated.

In(S,OH). Various compounds of In have been used, with some success,for buffer layers. In(OH)3 was grown on CIS (In-rich) films from a solution ofInCl3 with thiourea (which possibly acted to gradually increase pH rather than asa source of S) [30]. In spite of the higher blue response compared to a control CdScell [In(OH)3 is colorless as a film], the red response was poorer, leading to asomewhat reduced overall photocurrent. The fill factor was also less. A best effi-ciency of 9.5% was obtained, compared to 11.9% for the control using the samebatch of substrates. In the same paper, the deposition of In2S3 [based on later stud-ies, this may have been In(S,OH)] from a thioacetamide bath at a pH between 1and 2 was described, but with considerably poorer results (both photocurrent andfill factor were much lower).

Using essentially the same method for depositing In(S,OH) on CuInS2

films, the same group found slightly better performance (11.4%) with this bufferlayer than for CdS (10.8%), due to increase in photocurrent and also in open-cir-cuit voltage [31]. Since CuInS2 has a higher bandgap (1.5 eV) than CIS or CIGS,the fractional increase in photocurrent due to the improved blue response is larger.

These same In(S,OH) films were also investigated on CIGS substrates [32].XPS measurements (on CuGaSe2 used to prevent interference by In from the CIGS)showed that the deposited film was not a mixture of In2S3 and In(OH)3, as mightbe reasonably expected, but was a compound containing In, O, and S. Higherthioacetamide concentrations resulted in better device performance, while the Inconcentration was not found to influence the performance in any reproducible way.The completed cells (with ZnO window layer) required an anneal (2 min at 200°C)and light soak; the performance without this anneal was much poorer. The best cellgave an efficiency of 15.7%. In comparison with CdS buffers, the VOC increased,the ISC decreased, and there was no change in the fill factor. Overall, the efficien-cies were only slightly less than for cells using CdS. For CIGS substrates depositedon Corning “Pyrex-type” glass (good cells are usually deposited on soda-lime glassand are beneficially affected by Na� diffusion through the film), the In(S,OH) cells


were actually better than the CdS ones. In fact, in contrast to CdS cells, the effi-ciency of the In(S,OH) ones were independent of the type of glass used for the sub-strate. Capacitance–voltage measurements indicated a 10 times increase in accep-tor concentration of the absorber layer and narrowing of the space charge layer,compared to CdS. This explained the decrease observed in the red response (theblue response was better) and the increase in the VOC, and also the independenceon the glass, since acceptor doping occurred by the CD process rather than by Na�

diffusion. Since the CD process did not change the bulk of the CIGS, it was as-sumed that the doping changes were due to surface (and grain boundary) effects.

A list of cells made on CIS, CuInS2, and different CIGS substrates usingIn(OH,S) buffer layers is given by Hariskos et al. [32]. They note that the fill fac-tor of these cells drops with time when the cells are kept in the dark but that illu-mination (by light absorbed by the In(OH,S) film) reverses the degradation. Thisbehavior suggests that the degradation is due to adsorption of some species (oxy-gen, water?) that reacts with photogenerated electrons and/or holes.

Sn(O,S). In the same paper, films of SnO2 and Sn(O,S)2—the latter ap-parently a mixture of oxide and sulphide—were used as buffer layers on CIS andCIGS but with all cell parameters lower than when using CdS and efficiencieslower than those obtained using In(OH,S)

9.1.5 Other Cells and Related Studies

9.1.5.1 Other Heterojunctions and Devices

There are several studies on heterojunctions, other than those already described,formed with at least one CD semiconductor. These will be described briefly here.

The earliest of these studies was on PbS. PbS can have either p- or n-typeconductivity, although CD PbS is usually p-type. Based on the belief that the p-type conductivity may be due to alkali metal cations from the deposition solution,an alkali metal—free deposition, using lead acetate, thiourea, and hydrazine hy-drate was used [33]. While initially n-type, the film converted to p-type in air. At-tempts to stabilize the p-type material by adding trivalent cations to the depositionsolution were unsuccessful. However, deposition of the PbS on a trivalent metal,such as Al, did stabilize the n-PbS, at least for a time. In this way, p-n junctionswere made (the PbS close to the trivalent metal was n-type, while the rest of thefilm was p-type). Photovoltages up to 100 mV were obtained from these junctionsat room temperature and almost 300 mV at low temperatures (90 K).

PbS was also deposited on single-crystal n-and p-type Ge [34]. The PbS wasepitaxial (111) with the (111) Ge (Ge has a 5% smaller lattice spacing than PbS).A photovoltage was measured from the junctions. The photoresponse extended to1.75 �m for the junctions on p-type or intrinsic Ge and to 3.35 �m on n-type Ge.The difference could not be explained, although it can be noted that these values


correspond approximately to the bandgaps of Ge and PbS, respectively. Hot probemeasurements indicated that the PbS layers were n-type; however, it was men-tioned that this measurement may be affected by the junctions and that PbS chem-ically deposited onto glass by the same method normally gives p-type layers.

Investigations of junctions formed by CD of PbS on n-Si [35] and n-Sb2S3

on p-Si [36] and on p-Ge [37] have been made. In particular, the Sb2S3 junctionswere found to be much more PV active if a small amount of silicotungstic acid(STA) was added to the deposition solution. Conversion efficiencies of 5.2% on Siand ca. 4% on Ge could be calculated from the photocurrent–voltage characteris-tics. The STA resulted in formation of WO3 in the Sb2S3 film, and it was believedthat this may, at least in part, be responsible for the improvement. Of particular notewas the relatively large open-circuit voltage (nearly 0.7 V) obtained from the junc-tion with Ge; this value is almost as large as the Ge bandgap. It is tempting to won-der if this junction was not closer to a Schottky (metal–semiconductor) junction,where the Ge behaves as a metal, in which case, the maximum photovoltage is lim-ited by the Sb2S3 bandgap (ca. 1.7 eV). The p-Ge was highly doped (4 � 1018

cm�3), although less than might be expected if it were to behave as a metal.CuInSe2, deposited by CD (see Chap. 8), has been screened for photovoltaic

activity. In one study, CdS was evaporated on CIS that was chemically depositedonto conducting glass [38], while in the other the CIS was chemically depositedonto single-crystal Si [39]. The cells gave low activity, although the CdS/CIS cellgave a short-current photocurrent of 4 mAcm�2 (AM1 illumination), which isquite appreciable, if still low, and suggests further studies in this direction mightbe fruitful, in particular also using CD CdS.

CdO, a degenerate n-type semiconductor, was chemically deposited on sin-gle-crystal p-type Si [40]. The junction showed clear diode behavior, and, al-though no photovoltaic effect was observed, photocurrent was generated under re-verse bias. From the spectral response of the photocurrent, almost all of the currentgeneration occurred in the Si.

CdS/SnxS PV cells have been fabricated where the CdS was deposited byCD and the SnS deposited by a variant of CD where the substrate is dipped firstin a solution of one of the ions and then in the other without rinsing in between, aswould be the procedure for SILAR (see Sec. 2.11.1) [41]. While the cells showedvery low conversion efficiencies, the main emphasis was on Ag-doping of the CdSin order to increase the conductivity and the effect of this doping on the PV cells.An increase in efficiency from 0.03% to 0.08%, mainly as a result of an increasein short-circuit current, was obtained by doping the CdS with Ag. The doping wascarried out by an ion exchange process whereby the undoped CdS film was im-mersed in a solution containing Ag� complexed by thiosulphate.

A heterojunction between two different CD semiconductors—n-Ag2S de-posited on top of p-PbS—was fabricated [42]. Some photoconductivity was foundat wavelengths longer than that corresponding to the Ag2S bandgap (ca. 0.9 eV).


However, at shorter wavelengths, where light could be absorbed in the Ag2S, neg-ative photoconductivity (increase in resistivity with light) was found. This was ex-plained by a combined Dember effect/electrostatic attraction of electrons by Ag�

ions.Metal–CD semiconductor Schottky junctions have been examined as solar

cells. As for the p-n junctions described earlier, the addition of STA to both Sb2S3

and CdSe improved the cell parameters greatly [43]. AM1 efficiencies (for an-nealed films) of 7.2% and 5.5% were obtained using Pt contacts on annealed CdSeand Sb2S3, respectively. Again, the improvement due to the STA was attributed tothe presence of WO3. In this study, it was suggested that the WO3 might introducefavorable interface states in the devices. Whatever the reason for the improvement,the strong effect of the STA treatment and its applicability to two different semi-conductors (as well as to both solid-state and liquid junction cells) warrants furtherinvestigation. A similar study, using CD CdSe (with STA) deposited on poly(3-methylthiophene), the latter electropolymerized (or the polythiophene electrode-posited onto CdSe), was carried out [44]. Poly(3-methylthiophene) can be preparedeither p-semiconducting or doped to a metallic conductivity. While the undoped p-n junctions gave poor photoresponse, the Schottky-type doped thiophene-CdSejunction gave conversion efficiencies of 2.7%, which were stable for at least 72 hrof illumination. CdSe deposited without STA resulted in lower efficiencies (1.3%).

Thin-film transistors have been fabricated by depositing 50 nm of CdS ontoSiO2-covered n� Si and evaporating two A1 electrodes (source and drain) onto theCdS [45]. Similar devices were also made using CdS deposited on polyimide sub-strates with three (source, drain, and gate) evaporated metal electrodes and vari-ous sputtered insulator layers for the gate electrode.

9.1.5.2 Passivation Studies

Various sulphiding treatments have been known to passivate the surfaces of III–Vsemiconductors like GaAs and InP. Similar effects have been found with CD ofvery thin CdS films on InP surfaces. By deliberately oxidizing InP surfaces to pro-duce In and P oxides, it was shown that a standard CdS solution (only with a higher-than-usual concentration of ammonia) removed these oxides and prevented oxideregrowth [46]. The same treatment also removed P vacancies deliberately intro-duced by annealing in H2, as seen by photoluminescence studies, presumed due tofilling of P vacancies by S. Deposition of CdS (5–7 nm) onto InP both improvedthe C–V behavior and lowered the interface state density of SiO2/InP junctions byan order of magnitude. Further studies found that a pretreatment of the InP surfaceby a solution of NH4OH/thiourea improved the passivation but that CdS deposition(3–4 nm) was still important [47]. The NH4OH/thiourea treatment was believed toproduce a stable In-S terminated InP surface [48]. This latter study also reportedthat the SiO2/InP junction quality, measured mainly by the interface state density,was maximum at a very low CdS thickness (ca. 1 nm) and that this quality degraded


gradually as the CdS thickness increased, presumably due to interface states intro-duced by the thicker CdS. It appears that the NH4OH/thiourea treatment passivatesthe surface, while the very thin CdS protects it from reoxidation. This treatmentwas extended to various device structures based on ternary semiconductors(InA1As and InGaAs) with similar improvements due to removal of interface states[49]. These studies are clearly related to similar ones on CI(G)S.

CD ZnSe has also been demonstrated to passivate surface states, 0.92 eV be-low the conduction band edge (measured by thermally stimulated exoelectronemission) on single crystal GaAs. This passivation resulted in bandgap lumines-cence from the originally non-luminescent GaAs [49a].

9.2 PHOTOELECTROCHEMICAL CELLS (PECs)

9.2.1 Introduction and Background

As noted earlier, there are numerous studies on the photoelectrochemical (PEC)properties of CD films. Many, if not most, of these studies describe the prepara-tion of the films and some PEC properties. In such cases, rather than describe eachstudy separately, it is more useful and efficient to tabulate the results, providingimportant cell parameters together with the reference. Additional relevant infor-mation will be given separately for each individual reference. However, no at-tempt is made to cover all the individual studies in any detail, but rather to giveenough information to allow the reader to decide whether it may be worthwhile torefer to the original reference. Some studies that treat the PEC properties of thesefilms in a more fundamental way will be discussed separately. An important issueis whether the films have been annealed or not and under what conditions; an-nealed films usually give better performance (normally much better) than nonan-nealed ones.

FIG. 9.5 Schematic diagram of a PEC.


We start with a brief introduction to PECs. Figure 9.5 presents a schematicdiagram of a PEC showing a semiconductor film on a substrate—the photoelec-trode—connected through an external meter and/or load to a second electrode (thecounterelectrode). The two electrodes are immersed in an electrolyte, and the semi-conductor film is exposed to illumination. If the substrate is transparent, conduct-ing glass, the light can pass first through the substrate and then to the semiconduc-tor (and the glass can also function as the window of the cell); this configuration isknow as a backwall cell, in contrast to the normal frontwall cell (shown in Fig. 9.5),where light is incident directly on the nonsubstrate side of the semiconductor film.

The following discussion assumes that the semiconductor crystal size islarge enough so that charge transport is dominated by a space charge layer in thesemiconductor. This is typically the case when chalcogenide films have been an-nealed at temperatures of ca. 400°C or more, where the crystal size is typically ofthe order of hundreds of nanometers. This assumption is usually not valid for as-deposited CD films or for those annealed at low temperatures (e.g., 250°C or less).The mode of operation of such films is treated in Section 9.2.2.2.

Figure 9.6a shows the band diagram of a semiconductor–electrolyte junc-tion (with an n-type semiconductor; most studies described here have used filmsthat give n-type response). The band bending represents the situation in the darkat equilibrium (flat Fermi level) or under illumination at short circuit (i.e., the pho-toelectrode (or photoanode) is short-circuited to the counterelectrode). (We ignorecomplications of quasi-Fermi levels due to any nonequilibrium situation in thelight; it does not change the simple picture for our purposes). Photogenerated elec-trons (in the conduction band) and holes (in the valence band) are spatially sepa-rated by the space charge layer and flow in opposite directions. The holes flow tothe semiconductor/electrolyte junction and oxidize some species in the elec-trolyte. The electrons flow through the back (ohmic) contact to the external cir-cuit. Electrons thus flow to the counterelectrode and reduce some electrolytespecies. (For a p-type semiconductor, a photocathode, the direction of charge flowis opposite: Electrons flow to the semiconductor/electrolyte junction and holes tothe back contact.) The resulting current flowing is the short-circuit current (ISC).For a regenerative PEC, the electrolyte species that are oxidized and reduced forma single redox couple. A common example is the polysulphide redox solution. Thereactions occurring at the photoanode are:

S2� � 2h� → S; S � S2� → S22� etc. up to approximately S5

2� (9.1)

while at the counterelectrode, the reverse occurs:

S22� � 2e� → 2S2� or S � 2e� → S2� (9.2)

The result is no net change in the PEC (hence the term regenerative) and electric-ity is produced.

The regenerative PEC has been the type predominantly studied using CD


films. However, there is another important type, where the anodic and cathodic re-actions occur with different redox species. The holy grail for this type of cell hasbeen to photoelectrolyze water to hydrogen and oxygen. While this goal has at-tained only limited success, the search has led to very decided success in other, re-lated directions, usually connected with photo-oxidizing adsorbed layers of pollu-tants or bacteria on TiO2 [50].

If the photoanode is not connected to the counterelectrode, then currentcannot flow, and instead the potential of the photoanode (the Fermi level) risesuntil balanced by recombination of the photogenerated charges (Fig. 9.6b). Thedifference between this potential under illumination and the original potential(more correctly, the potential of the counterelectrode, which ideally is equal tothat of the nonilluminated photoelectrode) is the open-circuit voltage (VOC) of thecell.

FIG. 9.6 Band bending in a PEC (a) in the dark or under illumination at short circuit and(b) under illumination at open circuit.

A

B


In practice, in order to generate electrical power, the cell must operate un-der conditions where both current and voltage are generated, as with the photo-voltaic cells described earlier. This situation is shown in Figure 9.7, which givesthe photocurrent/photovoltage characteristics of the cell. The maximum power(Pmax) is generated when the load is such that the product of the current (IP) andvoltage (VP) is a maximum. The shape of the photocurrent/photovoltage charac-teristic, which determines Pmax, is quantified as the fill factor (FF), which is de-fined as the ratio between Pmax and the product of ISC and VOC, i.e.,

FF � �ISC

P�ma

Vx

OC� (9.3)

and is given either as a fraction or, commonly, as a percentage.An important feature of photoanodes is that the photogenerated holes, which

are normally very strongly oxidizing, may oxidize the semiconductor instead of, oras well as, the electrolyte species. This phenomenon is known as photocorrosion.For the purposes of the limited explanation of PECs given here, it is enough to notethat continuous photocorrosion will destroy the photoelectrode. (This need not nec-essarily occur if corrosion is confined to the semiconductor surface.)

The substrate on which the semiconducting photoelectrode is deposited isimportant, not just as an ohmic contact to extract charge (usually electrons, sincemost photoelectrodes are photoanodes), but also because the substrate should beelectrocatalytically as inactive as possible toward the electrolyte species. The rea-son for this is explained in Figure 9.8. The heavy solid line shows a typical pho-

FIG. 9.7 Current–voltage characteristic of a photoanode showing maximum powerpoint (Pmax). The fill factor is given by the product VP � IP divided by the product of VOC

and ISC.


toanode response; it is assumed that dark currents are negligible. We now considerthe dark (or under illumination—there will be no difference) current–voltagecharacteristic of any exposed substrate. This is shown for a poor electrocatalystand a better (still not good) electrocatalyst. The higher the electrocatalytic activ-ity of an electrode, the greater the current at a given bias. For a photoanode, thecurrent in the fourth quadrant is the most important (for a photocathode, it is thatin the second quadrant). If the positive current of the (no dark current) photoan-ode and the negative current of the exposed substrate are now added, the effect ofthe poor electrocatalyst is negligible while that of the better electrocatalyst is tostrongly reduce both the VOC (from VOC to VOC(1)) and the fill factor of the PEC.This is the reason that Ti, a poor electrocatalyst for most redox systems, is used socommonly as a substrate for photoanodes. Another substrate sometimes used is Cr(also a poor electrocatalyst in general), e.g., for Bi2S3 [51] and Sb2Se3 [52]. Thisis in addition to the fact that these metals tend to form satisfactory ohmic contactsto n-type semiconductors. Note that for nanocrystalline films where a spacecharge layer may not exist, the concept of ohmic or nonohmic contact is not nec-essarily the same as for bulk semiconductors (see Sec. 9.2.2.2).

FIG. 9.8 The effect of exposed substrate on a PEC. The thick solid line gives the light-induced current–voltage characteristic. The thin solid line gives the net current–voltagecharacteristic when the substrate has the electrocatalytic activity given by the broken line(fair electrocatalyst). If a substrate with poor electrocatalytic activity is used, there is littleeffect of the substrate on the photocharacteristics in the fourth quadrant.


9.2.2 CdSe

9.2.2.1 Annealed Films

CdSe has been the most extensively studied semiconductor for PEC purposes.This is due to its fairly favorable bandgap for solar cell use (1.73 eV), comparedto the higher-bandgap CdS, although there are many, usually less detailed, PECstudies on CdS as well. Cell details are given in Table 9.1, specific comments fol-low.

In Ref. 53, the main purpose of the study was to investigate PEC propertiesof CD CdSe films rather than to optimize the actual solar cell (e.g., CdSe anneal-ing was carried out at the relatively low temperature of 280°C and the film thick-ness was only ca. 0.9 �m), hence the cell parameters are not as high as they couldbe. The CdSe/polysulphide junction was characterized by a number of techniques.The effects of an surface layer of CdS, due to exchange of Se by S from the poly-sulphide, were considered.

Preliminary PEC results in Ref. 54 were previously described [60]. Thisstudy was directed to optimization of the PEC parameters. The deposition wasbased on the ammonia/selenosulphate bath. The Ti substrates were treated with asuspension of Cd(OH)2 and the deposition carried out in a sealed tube to preventloss of ammonia and thereby to improve reproducibility. Several layers (aboutfive) were deposited to give a total thickness of ca. 2.5 �m, which was found tobe optimum. The first layer was annealed at 500°C in air (to improve film adher-ence) and the final film at 550°C, also in air. The annealed films were etched (9MHCl) and treated with a ZnCl2 solution. The maximum efficiency was obtained at50 mWcm�2 illumination (6.8%). At higher illumination intensities, the effi-

TABLE 9.1 PEC Parameters of CD CdSe

Isc Voc IlluminationEfficiency (%) (mAcm�2) (V) FF Electrolyte (mWcm�2) Refs.

Annealed0.15 1.92 0.23 0.34 Polysulphide 100 536.3 15.3 0.66 ca. 0.42 Polysulphide 67 544.9a 16.0 0.66 0.46 Polysulphide 100 (AM1) 554.4 8.4 0.85 0.62 Polysulphide 100 5611.7 11.8 0.59 0.67 Selenosulphate 40 57, 58

Not annealed0.1 0.06 0.1 0.28 Polysulphide Not defined 59

a Higher efficiency (5.5%) was found for lower solar illumination (42.5 mWcm�2).


ciency dropped, due to a sublinear increase in ISC with illumination at higher in-tensities.

The effects of annealing and film thickness on the PEC properties were in-vestigated in Ref. 55. The optimum annealing conditions were 470°C for one hourin air. ISC in particular was found to drop strongly both at lower and higher an-nealing temperatures. ISC and VOC, and therefore efficiency, increased stronglywith film thickness up to ca. 1.5 �m and then more gradually up to ca. 2 �m, af-ter which no further change was observed. Such thicknesses result from a numberof successive depositions. Larger-area photoelectrodes of 18 cm2 were also made;the efficiency of these cells dropped to 3.2% compared with small cells, mainlydue to a drop in fill factor and ISC.

A triethanolamine/ammonia/selenosulphate bath was used in the experi-ments of Ref. 56. Three depositions were employed to give a final film thicknessof 4 �m. The films were annealed at 500°C in N2 and etched in dilute aqua regia.The efficiency at this stage in a polyselenide electrolyte was ca. 0.5%. This studyconcentrated on the effects of a treatment consisting of a 50% HCl etch, which re-sulted in a black matte surface (indicating a high surface roughness), followed bya dip in an acidified ZnCl2 solution (a Zn2� dip has been previously been usedbeneficially for CdSe photoelectrodes [61]). This surface treatment improved allcell parameters. Using Kelvin probe studies, it was shown that the surface poten-tial (measured against a Pt vibrating probe in nitrogen ambient) changed by �0.31V after the treatment, interpreted as an increase in band bending due to change inthe surface charge. This change was the same as the increase in VOC of the cell.Analysis of the dark current–voltage characteristics of the cell showed a decreasein both the reverse saturation current and the ideality factor after the surface treat-ment. That Zn was present in the samples was verified by separate XPS measure-ments [62]. A spectral response study of these films was carried out, with an em-phasis on subbandgap response, which might be related to surface states [63]. Thesurface treatment decreased the subbandgap response (although it should be re-marked that the subbandgap response of the untreated film was unusually high,extending out to 1 �m and with relatively high quantum yields for a subbandgapresponse). It also preferentially increased the short-wavelength response, inter-preted as a reduction in surface recombination. The beneficial effect of the surfacetreatment was found to occur not just in polyselenide electrolyte, but also in a fer-rocyanide electrolyte. This suggests that the effect of the Zn is not related to for-mation of ZnSe at the surface, as might be thought. On the other hand, a large partof the improvement in polyselenide electrolyte was due to the HCl matte etch, andit is not clear how much of the improvement in the ferrocyanide electrolyte wascaused by this part of the treatment.

In Ref. 57, a small concentration (� 10�5 M) of silicotungstic acid(H4SiW12O40) was added to a triethanolamine/ammonia/selenosulphate bath,


deposition carried out at 40°C and annealing in air at 430°C (the CdSe remainedcubic, in contrast to the hexagonal form usually obtained after annealing at suchtemperatures). A large increase in the PEC performance was obtained, comparedto photoelectrodes deposited without the silicotungstic acid. All cell parameterswere increased, but the major effect was on ISC. The stability of the PEC also im-proved: It was stable (at 40 mW-cm�2 illumination) for greater than 2800 C/cm2,although the selenosulphate electrolyte is unlikely to be stable in a PEC over thelong term. A follow-up study of these films [58] showed the presence of WO3 inthe films. From the dark current–voltage characteristics, the reverse saturationcurrent, ideality factor, and series resistance all decreased as a result of the modi-fication. Two possibilities were put forward to explain the effect of the silico-tungstic modification. One was that formation of a CdSe/WO3 heterojunction oc-curred, improving the charge transfer at the semiconductor/electrolyte interface.Another possibility, if a true heterojunction exists, is reduction in majority carrier(electron) transfer to the electrolyte, as in a metal/insulator/semiconductor device.In another study, it was suggested that the improved PEC response was due tocharge transfer catalysis by the W-containing groups adsorbed at the CdSe crys-tal surfaces [64]. It should also be noted that the silicotungstic acid modificationalso improved the performance of solid-state cells (see Sec. 9.1.5.1) where nocharge transfer to an electrolyte is involved, suggesting (although not proving)that electrocatalysis is not the reason for the improved PEC behavior.

9.2.2.2 Nonannealed Films

In Ref. 59, films 0.8 �m thick were deposited from an ammonia/selenosulphatebath. Various configurations of PECs were studied.

A study of the mode of operation of nanocrystalline CdSe photoelectrodeswas carried out [65]. This was based on the expectation that nanocrystals, being(usually) much smaller than typical space charge layer widths, would not sup-port such a space charge layer and therefore that some other mechanism forcharge separation should be considered. It was originally reported that nanocrys-talline CD CdS [66] and CdSe [67] photoelectrodes, which normally gave pho-tocurrent–voltage behavior characteristic of an n-type semiconductor, gave p-type behavior after etching in dilute HCl, an example of which is shown inFigure 9.9. The study by Kronik et al. [65] used surface photovoltage (SPV)spectroscopy and X-ray photoelectron spectroscopy (XPS) to investigate this ef-fect. It was concluded that the CdSe itself was close to intrinsic, as might be ex-pected for very small nanocrystals [the crystal size in this study was 4–5 nm; nosuch apparent reversal of semiconductor type was found for considerably larger(ca. 20 nm) crystal size]. The direction of photocurrent flow, rather than beingdetermined by an electric field (the space charge layer) in the semiconductor,was determined by trapping of charge carriers at the individual nanocrystal sur-faces. Both electrons and holes could be trapped at the surface; however, the


charge that was preferentially trapped (longer lifetime of trapped charge, whichusually correlated with deeper trap sites) would preferentially be transferred tothe electrolyte. Such a scenario is reasonable due both to the longer residencetime of the charge at the surface as well as better overlap of a deeply trappedcarrier with the redox species, which, for the polysulphide system used in theseexperiments, is located deep in the bandgap. The presence of such charge trapsand the effect of water on them was further substantiated in later studies usingphotoluminescence [68,69] and scanning probe spectroscopies [70]. See Section4.2.7.1 for a description of surface trapping in CD CdSe measured by variousphotoluminescence studies.

The photovoltage of such a cell comes not from neutralization of a built-inspace charge layer, but from change in the Fermi level in the almost intrinsicnanocrystal upon illumination. This change in Fermi level is also determined bythe relative trapping of electrons and holes. If no trapping occurred and no chargeextraction took place, then illumination would not change the potential of thesemiconductor appreciably, since both electron and hole Fermi levels wouldchange by (close to) the same amount and in opposite directions. However, to takea simple example, if the holes are deeply trapped and the electrons are not trappedat all, then the shift in the overall Fermi level is determined almost completely by

FIG. 9.9 Current–voltage characteristics under chopped illumination (� AM1) of ananocrystalline (4–5 nm) CdSe film, deposited by CD, in a polysulphide electrolyte. Thetwo characteristics are for as-deposited CdSe (top) and after etching with dilute HCl (bot-tom). (After Ref. 67.)


the electron concentration; i.e., the Fermi level rises on illumination, similar towhat happens in an n-type semiconductor. The opposite will occur if electrons aretrapped and holes are free. In practice, both charges are usually trapped to a greateror lesser extent, and the shift in Fermi level, and therefore the photovoltage direc-tion and magnitude, will be dependent mainly on the relative trapping depths ofthe two carriers.

While it has not been studied, it is probable that there exists a range of crys-tal sizes, possibly quite wide, where both surface trapping and space charge layereffects contribute to the PEC functioning of the photoelectrode.

While the cell efficiencies of these films were not specifically investigated,best parameters of 2 mAcm�2 (ca. AM1 illumination; quantum efficiencies in-creased with decreasing illumination intensity due to diffusion limitations in thenanoporous film); 0.5 V and ca. 50% fill factor were obtained. However, greatvariation in these parameters were obtained; one reason for this can be seen froma consideration of Figure 9.9. If a CdSe film is etched, but less than optimally(shorter time, more dilute HCl), it is clear that after a certain, unique etch treat-ment, zero net photocurrent will be obtained. The actual photocurrent (and otheroutput parameters) of the film is a balance between photoanodic and photoca-thodic currents.

A word on “ohmic” contacts to space charge layer–free nanocrystallinefilms. The ability of a contact to function as an ohmic contact to such films is de-termined by the offset between the metal contact Fermi level and the semicon-ductor energy bands, rather than by a potential barrier in the form of an electricfield in the semiconductor. It seems that, in practice, most conductors act as goodsinks for photogenerated charges in nanocrystalline semiconductors. One reasonfor this may be that accumulation of one type of charge (the other is usually re-moved rapidly by the electrolyte at the very high contact area between nanocrys-tals and electrolyte) will charge the particle and raise the energy levels until theaccumulated charge can flow to the sink.

9.2.3 CdS

In contrast to CdSe, most studies on CdS involved either nonannealed films orfilms annealed at relatively low temperatures. Relative efficiencies (i.e., takinginto account that the efficiency of the higher-bandgap CdS will be lower than thatof CdSe due to its lower light absorption) are therefore low. Additionally, mostPEC studies on CD CdS involve doped films, where ions of the dopant were addedto the deposition solution.

In Ref. 71, Al doping (0.1 wt% Al added to the deposition bath; not neces-sarily the same percentage in the films) improved the PEC properties. Annealing(in H2) at 200°C, which removes S and should make the films more n-type, im-proved the cell parameters of both doped and undoped films [72,76]. Stability data


in polysulphide electrolytes were presented. In doping (optimum 0.01% In) alsoimproved the PEC characteristics [73].

The study in Ref. 74 followed essentially the same deposition procedure asbefore, but on Cr-plated steel (compared with steel in the previous studies; seeSec. 9.2.1 for an explanation of the effect of Cr plating based on the poor electro-catalytic activity of Cr) and with a CuCl:KCN etch at 90°C (the reason for thisspecific etch was not explained). These modifications led to improved PEC re-sponse, which was further improved using As-doped CdS, deposited by addingAsCl3 to the deposition bath.

In general, there is an optimum doping concentration, which varied fromdopant to dopant. The dopants studied here were donors (increased n-type CdS).Increase in doping density can increase all parameters; but if too much dopant ispresent, the parameters can degrade due to a narrower space charge layer (poorerISC due to decreased collection efficiency in the red) and increased recombinationdue to impurity centers. For nanocrystalline nonannealed (or low-temperature-an-nealed) films, where there may be no space charge layer as discussed previously,the effect of recombination centers will still be valid—maybe even more impor-tant due to the lack of a space charge layer to separate charges.

Cu doping of CdS has been investigated [77]. Since the light intensity wasnot specified, the cell parameters are not given in Table 9.2. The doping caused asmall increase in ISC and an equally small decrease in VOC, with no appreciablechange in efficiency, although it is arguable if these changes are significant. Theseelectrodes were also used with a Ag2S storage electrode in a photoelectrochemi-cal storage cell.

Photoelectrochemical characterization was also carried out on CdS films us-ing different sizes of CdS nanocrystals [75]. VOC increased with decreasing crys-tal size from 0.58V (75 nm) to 0.68 V (5 nm). Surprisingly, ISC was not dependent

TABLE 9.2 PEC Parameters of CD CdS

Isc Voc IlluminationEfficiency (%) (mAcm�2) (V) FF Electrolyte (mWcm�2) Ref.

0.003 0.06 ca. 0.11 0.33 Polysulphide 100 710.0075 (Al-doped) 0.125 ca. 0.11 0.30 Polysulphide 1000.04 (Al-annealed) 0.5 0.24 0.35 Polysulphide 100 720.01 0.17 0.17 0.44 Polysulphide 100 730.05 (In-doped) 0.4 0.23 0.53 Polysulphide 100 730.022 0.3 0.175 0.42 Polysulphide 100 740.05 (As-doped) 0.5 0.22 0.44 Polysulphide 100 74— ca. 0.45 0.58–0.6 Polysulphide 60 75

or sulphite


on size, as would be expected due to the increasing bandgap and therefore low-ered light absorption with decreasing crystal size.

Besides evaluating photoelectrodes for use in PECs, photoelectrochemicalcharacterization can be used for other purposes. For example, photocurrent spec-tra of CD CdS has been used to measure the semiconductor bandgap (as I2 vs. h�),and agreement between the bandgap values measured by this method and by ab-sorption spectroscopy for as-deposited and annealed films was found [78].

9.2.4 Other Photoelectrodes

A number of other CD semiconducting materials as photoelectrodes have been re-ported, the basic PEC characteristics are given in Table 9.3. Further details can befound in the original references.

Most semiconductors described in this chapter gave n-type response (as al-ready shown, for nanocrystalline semiconductors this does not necessarily meanthat the semiconductors are actually n-type but rather that the net photogeneratedhole current is to the electrolyte while that of the electrons is to the substrate). Oneexample of a “p-type” photoelectrode is nanocrystalline CD PbSe [83]. Figure 9.10shows the response of such a film. The crystal size of the PbSe in this film is ca. 4nm, with regions of crystals two to three times larger (the crystal size distributionis bimodal—see Ref. 88). Interestingly, a single crystal of PbSe was not found togive any photoresponse at all under the same conditions, and the current–voltagebehavior of the single crystal was essentially ohmic rather than the clearly asym-

FIG. 9.10 Chopped illumination (100 mW-cm2) current–voltage characteristics of ananocrystalline PbSe film deposited from a citrate/selenosulphate bath at 60°C. The elec-trolyte is the original solution from which the film was deposited. (After Ref. 83.)


TABLE 9.3 PEC Parameters of Miscellaneous Semiconductor Films

Isc Voc IlluminationEfficiency (%) (mAcm�2) (V) FF Electrolyte (mWcm�2) Ref.

Sb2S3

0.8 (as-dep.) 2.0 0.44 0.37 Polyiodide 40 792.0 (ann.)a 4.0 0.48 0.443.9 (�STA)b 5.6 0.54 0.520.008c 0.14 0.155 0.41 Ferrocened 100 80

Sb2Se3 (nonannealed)0.06 0.45 0.37 ca. 0.3 Polyiodide 80 (IR filter) 52

Bi2S3 (nonannealed)0.12 0.34 0.39 0.55 Polysulphide 60 (water filter) 51

Ag2S (nonannealed)–– 0.1 0.1 –– Polysulphide 100 810.088 0.25 0.09 ca. 0.25 Polysulphide 200 82

PbSe (nonannealed)p-type response

0.029 1.8 0.08 0.4 Pb2�/SeSO32�e

100 83

SnS2 (nonannealed)0.006 0.019 0.45 0.43 KCl 100 84

(Cd, Pb) S0.14 ann.f 1.4 0.2 0.39 Polysulphide 75 850.83g no ann. ca. 2.8 ca. 0.5 0.3 Polysulphide 50 86

(Cd, Zn) S(nonannealed)h

0.04 0.065 0.31 0.43 Polysulphide 22 87

(Cd, Hg)S (annealedat 320�C in air)i

0.36 1.6 0.51 0.33 Polysulphide 75 85

a Annealed at 300�C in N2.b With silicotungstic acid and annealed at 300�C in N2.c As deposited.d In dimethylsulphoxide solution.e The solution from which the PbSe was deposited.f Annealed at 320�C in air; 0.11 mole fraction Pb––efficiency decreased as Pb increased.g For Cd0.925Pb0.075S.h For Cd0.7Zn0.3S.i For Cd0.82Hg0.18S.


metric diode behavior of the nanocrystalline film. Note also the p-type response ofetched CdSe (also CdS) nanocrystalline photoelectrodes discussed in Section9.2.2.2. Photoresponse spectra of PbSe films of different crystal sizes, reflectingthe varying bandgap due to size quantization, are shown in Figure 10.6.

Some ternary compounds have also been used as photoelectrodes, with op-timum efficiencies reported in Table 9.3. For Cd1�xPbxS, the efficiency increasedfrom 0.3% (pure CdS) to 0.83% (x � 0.075) and then decreased sharply to give aneven lower efficiency than for pure CdS (at x � 0.1), continuing to drop moreslowly at higher Pb fractions [86], although another study on this material foundthe efficiency of all alloys to be less than that of the pure CdS [85]. A similar trendof initial increase in efficiency was found for Cd1�xZnxS, with an increase in ef-ficiency from 0.01% (CdS) to a maximum of 0.04 (Cd0.7Zn0.3S), dropping to theCdS value at x � 0.4 and continuing to drop slowly at higher Zn values [87]. Theinitial increase in efficiency for the (Cd,Pb)S film with increasing Pb content wasprobably due to the lower bandgap and therefore increased absorption. The reasonfor the subsequent decrease in efficiency with further increase in Pb cannot be ex-plained so simply, although it may be noted that CD PbS is not very photoactive,possibly due to its relatively high conductivity. For the (Cd,Zn)S electrodes, sincethe bandgap of ZnS is much higher than that of CdS, the initial increase in effi-ciency (ISC also increases) is unexpected, although a possible reason is better dop-ing characteristics, resulting in an optimum resistivity; The series resistance ofthese photoelectrodes was a minimum at x � 0.3. For (Cd,Hg)S films, while thebandgap decreased with addition of Hg, the efficiency of films with a little (0.04mole fraction) Hg was much lower than those of pure CdS but increased with fur-ther addition of Hg until efficiencies close to those of pure CdS were reached at0.18 mole fraction Hg [85].

9.2.5 Coupled Photoelectrodes

Some studies involving coupled photoelectrodes of two or even three semicon-ductors, with at least one deposited by CD, have been reported. A relatively pop-ular subject is sensitization of nanocrystalline wide-bandgap semiconductors withlower-bandgap semiconductors. This is due to the wide interest in dye-sensitizedTiO2 solar cells [89]. A range of semiconductors has been deposited on wide-bandgap semiconductors using a variety of techniques, most commonly by dip-ping the wide-bandgap semiconductor film (usually TiO2) into solutions of ametal salt and then a chalcogenide solution, but also using electrodeposition andchemical vapor deposition (Refs. 90 and 91 give some examples). Charge separa-tion efficiency often is improved compared to the absorbing semiconductor itself,usually due to electron injection into the conduction band of the large-bandgapsemiconductor, which reduces electron–hole recombination. The stability of thesemiconductor is also often improved.


In such a coupled system, CdSe has been chemically deposited onto TiO2

(the latter prepared by both spray-painting and by screen printing) from a tri-ethanolamine/ammonia/selenosulphate bath [92]. The films were annealed at400°C in air. The quantum efficiencies of such films using spray-painted TiO2 andpolysulphide electrolyte, were found to be ca. 20 times higher than films of CdSeby itself (maximum quantum efficiency was reported to be over 0.6, and light-to-electricity conversion was 1%). Much lower values were found for the CdSe-sen-sitized (lower-surface-area) screen-printed TiO2 films, for nonannealed films, andin ferro/ferricyanide electrolyte. The spectral response of the CdSe/screen-printedTiO2 electrode exhibited a pronounced response beyond the bulk bandgap ofCdSe. This subbandgap response was absent in the CdSe electrode. It also was notapparent in the absorption spectrum of the coupled electrode, although this ab-sorption spectrum was not clearly defined (it was obtained from diffuse re-flectance measurements).

CdSe was deposited on CdS (both deposited by CD) and subjected to dif-ferent annealing temperatures [93,94]. The purpose was to see if the CdS/CdSeheterojunction affected the PEC properties. The main effect of the coupled systemcompared to only CdSe was to improve the stability of the photoelectrode inferro/ferricyanide electrolyte (a partially stabilizing electrolyte for CdSe). Thespectral response of the coupled system (measured, as usual, at low light intensi-ties) was closer to pure CdS than to CdSe, although the values of ISC were similarfor both at solar intensities, indicating an illumination-dependent behavior. Depo-sition of ZnO (by dipping in ZnAc2/methanol and annealing at 400°C) improvedthe stability of the system greatly, although with a decrease in ISC, so it is difficultto know how much of the increase in stability was due to the lower ISC and howmuch to other factors (stability normally will increase at lower currents). Changesin surface morphology, related to formation of Cd2Fe(CN)6, were measured on allsamples. Various hypotheses were put forward to explain the effects of the cou-pled system and of the ZnO on the stability.

REFERENCES1. RL Call, NK Jaber, K Seshan, JR Whyte Jr. Sol. Energy Mater. 2:373, 1980.2. E Fatas, T Garcia, C Montemayor, A Medina, E Garcia Camarero, F Arjona. Mater.

Chem. Phys. 12:121, 1985.3. TL Chu, SS Chu, C Ferekides, CQ Wu, J Britt, C Wang. J. Appl. Phys. 70:7608, 1991.4. TL Chu, SS Chu, N Schultz, C Wang, CQ Wu. J. Electrochem. Soc. 139:2443,

1992.5. ME Özsan, DR Johnson, M Sadeghi, D Sivapathasundaram, D Lincot, B Mokili, M

Froment, J Vedel, LM Peter, G Goodlet, RC Walker. In: 13th ECPV Solar EnergyConf., Nice, France, 1995, p 2115.

6. PN Gibson, ME Özsan, JM Sherborne, D Lincot, P Cowache. In: 2nd World Conf. &Exhib. on PV Sol. Energy Conversion. Vienna, Austria, 1998, p 1089.


7. J Tousková, D Kindl, L Dobiásová, J Tousek. Sol. Energy Mater. Sol. Cells 53:177,1998.

8. D Lincot, A Kampmann, B Mokili, J Vedel, R Cortes, M Froment. Appl. Phys. Lett.67:2355, 1995.

9. A Kampmann, P Cowache, B Mokili, R Ortega-Borges, D Lincot, J Vedel. In: 12thECPV Solar Energy Conf., Amsterdam, the Netherlands 1994, p 664.

10. L Kronik, L Burstein, M Leibovitch, Y Shapira, D Gal, E Moons, J Beier, G Hodes,D Cahen, D Hariskos, R Klenk, HW Schock. Appl. Phys. Lett. 67:1405, 1995.

11. D Hariskos, M Powalla, N Chevaldonnet, D Lincot, A Schindler, B Dimmler. ThinSolid Films 387:179, 2001.

12. D Lincot, R Ortega-Borges, J Vedel, M Ruckh, J Kessler, KO Velthaus, D Hariskos,H-W Schock. In: 11th European PVSEC. Montreux, Switzerland, 1992, p. 870.

13. J Kessler, KO Velthaus, M Ruckh, R Laichinger, HW Schock, D Lincot, R Ortega, JVedel. In: 6th Int. PV Sol. Energy Conf., New Delhi, India, 1992, p 1005.

14. K Ramanathan, RN Bhattcharya, J Granata, J Webb, D Niles, MA Contreras, H Wies-ner, FS Hasoon, R Noufi. Proc. 26th IEEE PVSC. Anaheim, CA, 1997, p 319.

15. K Ramanathan, H Wiesner, S Asher, D Niles, RN Bhattcharya, J Keane, MA Contr-eras, R Noufi. In: 2nd World Conf. & Exhib. on PV Sol. Energy Conversion: Vienna,Austria, 1998, p 477.

16. T Wada, S Hayashi, Y Hashimoto, S Nishiwaki, T Sato, Z Negami, M Nishitani. In:2nd World Conf & Exhib. on PV Sol. Energy Conversion. Vienna, Austria, 1998, p403.

17. A Kylner. J. Electrochem. Soc. 146:1816, 1999.18. T Nakada, A Kunioka. Appl. Phys. Lett. 74:2444, 1999.19. T Nakada. Thin Solid Films 361:346, 2000.20. A Kylner. J. Appl. Phys. 85:6858, 1999.21. L Kronik, U Rau, JF Guillemoles, D Braunger, HW Schock, D Cahen. Thin Solid

Films 361:353, 2000.22. MJ Furlong, D Lincot, M Froment, R Cortés, AN Tiwari, M Krejci, H Zogg. In: 14th

ECPV Sol. Energy Conf., Barcelona, Spain, 1997, p 1291.23. V Nadenau, D Hariskos, HW Schock, M Krejci, FJ Haug, AN Tiwari, H Zogg, G

Kostorz. J. Appl. Phys. 85:534, 1999.24. R Ortega Borges, D Lincot, J Vedel. In: 11th ECPV Sol. Energy Conf., Montreux,

Switzerland, 1992, p 862.25. K Kushiya, T Nii, I Sugiyama, Y Sato, Y Inamori, H Takeshita. Jpn. J. Appl. Phys.

35:4383, 1996.26. T Nakada, K Furumi, K Sonoda, A Kunioka. In: 2nd World Conf. & Exhib. on PV

Sol. Energy Conversion, Vienna, Austria, 1998, p 1173.27. A Ennaoui, M Weber, M Saad, W Harneit, MC Lux-Steiner, F Karg. Thin Solid Films

361:450, 2000.28. A Ennaoui, M Weber, CD Lokhande, R Scheer, HJ Lewerenz. In: 2nd World Conf.

& Exhib. on PV Sol. Energy Conversion, Vienna, Austria, 1998, p 628.29. A Ennaoui, M Weber, R Scheer, HJ Lewerenz. Sol. Energy Mater. Sol. Cells 54:277,

1998.30. KO Velthaus, J Kessler, M Ruckh, D Hariskos, D Schmid, HW Schock. In: 11th

ECPV Sol. Energy Conf., Montreux, Switzerland, 1992, p 842.


31. D Braunger, D Hariskos, T Walter, HW Schock. Sol. Energy Mater. Sol. Cells 40:97,1996.

32. D Hariskos, M Ruckh, U Ruhle, T Walter, HW Schock, J Hedstrom, L Stolt. Sol. En-ergy Mater. Sol. Cells 41-2:345, 1996.

33. J Bloem. Appl. Sci. Res. Section B 6:92, 1956.34. JL Davis, MK Norr. J. Appl. Phys. 37:1670, 1966.35. H Sigmund, K Berchtold. Phys. Stat. Sol. 20:255, 1967.36. O Savadogo, KC Mandal. Appl. Phys. Lett. 63:228, 1993.37. O Savadogo, KC Mandal. J. Phys. D Appl. Phys. 27:1070, 1994.38. KR Murali. Thin Solid Films 167:L19, 1988.39. JC Garg, RP Sharma, KC Sharma. Thin Solid Films 164:269, 1988.40. M Ortega, G Santana, A Morales-Acevedo. Solid State Electron. 44:1765, 2000.41. M Ristova, M Ristov. Sol. Energy Mater. Sol. Cells, 53:95, 1998.42. MJ Chockalingam, KN Rao, N Rangarajan, CV Suryanarayana. Phys. Status Solidi

(a) 2:K17, 1970.43. O Savadogo, KC Mandal. J. Electrochem. Soc. 141:2871, 1994.44. C Sene, HN Cong, M Dieng, P Chartier. Mater. Res. Bull. 35:1541, 2000.45. JH Schön, O Schenker, B Batlogg. Thin Solid Films 385:271, 2001.46. K Vaccaro, HM Dauplaise, A Davis, SM Spaziani, JP Lorenzo. Appl. Phys. Lett.

67:527, 1995.47. HM Dauplaise, K Vaccaro, A Davis, GO Ramseyer, JP Lorenzo. J. Appl. Phys.

80:2873, 1996.48. A Davis, K Vaccaro, HM Dauplaise, WD Waters, JP Lorenzo. J. Electrochem. Soc.

146:1046, 1999.49. K Vaccaro, A Davis, HM Dauplaise, SM Spaziani, EA Martin, JP Lorenzo. J. Elec-

tron Mater. 25:603, 1996.50. A Fujishima, K Hashimoto, T Watanabe. TiO2 Photocatalysis: Fundamentals and Ap-

plications; BKC Inc., Tokyo: 1999.51. RN Bhattacharya, P Pramanik. J. Electrochem. Soc. 129:332, 1982.52. RN Bhattacharya, P Pramanik. Sol. Energy Mater. 6:317, 1982.53. K Rajeshwar, L Thompson, P Singh, RC Kainthla, KL Chopra. J. Electrochem. Soc.

128:1744, 1981.54. RA Boudreau, RD Rauh. J. Electrochem. Soc. 130:513, 1983.55. RC Kainthla, JF McCann, D Haneman. Sol. Energy Mater. 7:491, 1983.56. KC Mandal. J. Mater. Sci. Lett. 9:1203, 1990.57. KC Mandal, O Savadogo. J. Mater. Sci. Lett. 10:1446, 1991.58. O Savadogo, KC Mandal. Mater. Chem. Phys. 31:301, 1992.59. SS Dhumure, CD Lokhande. Sol. Energy Mater. Sol. Cells. 29:183, 1993.60. DR Pratt, ME Langmuir, RA Boudreau, RD Rauh. J. Electrochem. Soc. 128:1627,

1981.61. G Hodes, D Cahen, J Manassen, M David. J. Electrochem. Soc. 127:2252, 1980.62. KC Mandal, O Savadogo. J. Mater. Sci. 27:4355, 1992.63. KC Mandal, KSV Santhanam. J. Mater. Sci. 26:3905, 1991.64. M Froment, H Cachet, H Essaaidi, G Maurin, R Cortes. Pure Appl. Chem. 69:77, 1997.65. L Kronik, N Bachrach-Ashkenasy, M Leibovitch, E Fefer, Y Shapira, S Gorer, G

Hodes. J. Electrochem. Soc. 145:1748, 1998.


66. G Hodes, A Albu-Yaron. Proc. Electrochem. Soc. 88-14:298, 1988.67. G Hodes, IDJ Howell, LM Peter. In: Photochemical and Photoelectrochemical Con-

version and Storage of Solar Energy. Tian, Z. W., Cao, Y., eds. Int. Acad. Publishers,Beijing, China, 1993, p 331.

68. E Lifshitz, I Dag, I Litvin, G Hodes, S Gorer, R Reisfeld, M Zelner, H Minti. Chem.Phys. Lett. 188, 1998.

69. E Lifshitz, I Dag, I Litvin, G Hodes. J. Phys. Chem. B 102:9245, 1998.70. B Alperson, I Rubinstein, G Hodes. Phys. Rev. B 6308:1303, 2001.71. CD Lokhande, SH Pawar. Mat. Res. Bull. 18:1295, 1983.72. CD Lokhande, SH Pawar. Sol. State Commun. 49:765, 1984.73. SH Pawar, LP Deshmukh. Mater. Chem. Phys. 10:83, 1984.74. LP Deshmukh, AB Palwe, VS Sawant. Solar Cells 28:1, 1990.75. KK Nanda, SN Sarangi, S Mohanty, SN Sahu. Thin Solid Films 322:21, 1998.76. CD Lokhande, MD Uplane, SH Pawar. Ind. J. Pure & Appl. Phys. 21:78, 1983.77. CD Lokhande, MD Uplane, SH Pawar. Sol. State Commun. 43:623, 1982.78. ME Özsan, DR Johnson, M Sadeghi, D Sivapathasundaram, G Goodlet, MJ Furlong,

LM Peter, A Shingleton. J. Mater. Sci. Mater. Electron. 7:119, 1996.79. O Savadogo, KC Mandal. J. Electrochem. Soc. 139:L16, 1992.80. LP Deshmukh, SG Holikatti, BP Rane, BM More, PP Hankare. J. Electrochem. Soc.

141:1779, 1994.81. SS Dhumure, CD Lokhande. Mater. Chem. Phys. 28:141, 1991.82. SS Dhumure, CD Lokhande. Sol. Energy Mater. Sol. Cells 28:159, 1992.83. S Gorer. PhD dissertation, Weizmann Institute of Science, Rehovot, Israel, 1996.84. CD Lokhande. J. Phys. D: Appl. Phys. 23:1703, 1990.85. M Skyllas-Kazacos, JF McCann, R Arruzza. Appl. Surf. Sci. 22/23:1091, 1985.86. LP Deshmukh, BM More, CB Rotti, GS Shahane. Mater. Chem. Phys. 45:145, 1996.87. LP Deshmukh, CB Rotti, KM Garadkar. Mater. Chem. Phys. 50:45, 1997.88. S Gorer, A Albu-Yaron, G Hodes. J. Phys. Chem. 99:16442, 1995.89. B O’Regan, M Gratzel. Nature 353:737, 1991.90. D Liu, PV Kamat. J. Phys. Chem. 97:10769, 1993.91. R Vogel, P Hoyer, H Weller. J. Phys. Chem. 98:3183, 1994.92. ME Rincon, O Gomez-Daza, C Corripio, A Orihuela. Thin Solid Films 389:91, 2001.93. ME Rincon, M Sanchez, A Olea, I Ayala, PK Nair. Sol. Energy Mater. Sol. Cells

52:399, 1998.94. ME Rincon, M Sanchez, J Ruiz-Garcia. J. Electrochem. Soc. 145:3535, 1998.


10Nanocrystallinity and SizeQuantization in ChemicalDeposited Semiconductor Films

Chemical deposition is a low-temperature technique compared with most othersemiconductor film deposition methods. This has both advantages and disadvan-tages. An obvious advantage is the simple processing often inherent in a low-tem-perature technique. What may be a more important advantage, however, is thatlow-temperature deposition techniques usually (although not always) result insmall crystal size. As recently as a decade ago, this would have been considered adecided disadvantage—large crystal size was almost always desired then, e.g., forphotovoltaic cells in order to minimize grain boundary recombination. However,with the increasing emphasis on nanostructured materials over the past decade,this characteristic of CD films is now often considered an advantage.

This chapter deals mainly with quantum size effects in CD nanocrystallinefilms. However, another, quite separate property of such films is related to thelarge percentage of atoms located on the surface of the nanocrystals of these films,e.g. �50% for a crystal size of a few nm; this is the effect of adsorption of molec-ular and ionic species on the nanocrystal surfaces. This aspect has been dealt withmuch less than has size quantization; therefore, it constitutes only a very small partof this chapter, mainly Section 10.2.3, which discusses the effect of adsorbed wa-ter on CD CdSe films. Section 9.2.2.2 deals in somewhat more detail with this par-ticular issue.


10.1 THE QUANTUM SIZE EFFECT:BACKGROUND

Quantum size effects in semiconductor nanocrystals became an important field ofresearch in the 1980s, when a number of groups, notably those of Brus at BellLabs and Henglein at the Hahn Meitner Institute, published seminal papers on theeffects of the size of semiconductor colloids on their optical properties and corre-lated crystal size with changes in electronic band structure.

Quantum size effects in semiconductor nanocrystals have been seen before,although the effect presumably was not realized. Early references to precipitatesformed when alkali, cyanide-containing selenide solution was added to ammoni-acal, cyanide-containing Cd2� described them as orange-yellow when precipi-tated in the cold and changing to red-brown when heated, and also noted thatfinely divided Cd-selenide varies in color from yellow to red-brown [1–3]. As isdescribed later, these color changes from normal dark-brown or black CdSe arethe most obvious and visual manifestations of the quantum size effect (or sizequantization).

The terms nanocrystals and quantum dots are often used interchangeably.Quantum dots, as used here, are invariably nanocrystals (amorphous materialscould, in principle, also exhibit quantum size effects as long as some electronicseparation between different particles occurs) that show quantum effects, whilenanocrystals may or may not be small enough to exhibit such effects.

Three-dimensional size quantization is due to localization of electrons andholes in a confined volume—e.g., a semiconductor nanocrystal—resulting in achange of the energy band structure. As the crystal size decreases below a certainlimiting size, associated with its exciton Bohr diameter, the spacing between lev-els in the bands becomes larger [the energy structure changes from a quasi-con-tinuum (band) to discrete, quantized levels] and the bandgap increases. This latterchange is manifested as a blue shift in the optical spectrum of semiconductorquantum dots (the “quantum size effect”). A simple way to visualize this effect isto consider a silicon atom (silicon being the best-known and most common semi-conductor), with two electrons in the 3s level and two in the 3p levels (Fig. 10.1).sp3 hybridization of the single s and three p levels leads to formation of four de-generate (i.e., equal-energy) sp3 levels, each containing one electron. Interactionof these levels with neighboring atoms results in splitting into bonding (�) and an-tibonding (�*) orbitals, with all electrons in the bonding orbitals and none in theantibonding ones. Up to here, the situation in Figure 10.1 is shown for a singleatom. However, as additional atoms are added, each atom contributes its own or-bitals and the localized orbitals of the single atom are gradually broadened into arange of molecular orbitals and, eventually, when the number of atoms becomesvery large, into the familiar filled valence band and empty conduction band char-acteristic of a semiconductor. If we consider the process from the other direction


(right to left), then the bandgap increases and the levels within the bands, whichin the bulk semiconductor are extremely (almost infinitesimally) close to neigh-boring levels, open up into discrete levels, as described previously.

This picture is reasonably valid for covalent silicon but rather simplistic formany of the semiconductors common in CD, which are usually mixed covalentand ionic. However, it serves to give a feeling for size quantization. For thosereaders who would prefer a more realistic interpretation for semiconductors withconsiderable ionic character, it is suggested that they construct a similar schemefor purely ionic materials and then “imagine” the required combination of ionicand covalent character.

There are many theoretical models to correlate the increase in semiconduc-tor bandgap with crystal size. However, for our purposes we will show only theoriginal model, known as the effective mass model, since this is the easiest to un-derstand, in spite of its limited accuracy.

The effective mass model is based on the energy of the lowest-energy par-ticle-in-a-box configuration, taking into account that the relevant mass term isgiven by a reduced effective mass, �, where � is given by

��1

� � �m1

e� � �

m1

h� (10.1)

The effective masses of electrons (me) and holes (mh) represent the masses thatthese charges appear to have when moving in the solid rather than in free space,and these vary from material to material. (In the size quantized regime, they canalso vary with crystal size, particularly for small quantum dots, hence the limita-tions of the effective mass model).

FIG. 10.1 Scheme showing (from left to right) how the relevant energy levels of siliconhybridize, interact with other atoms, split in a cluster, and eventually broaden into bands.(Adapted from Fig. 4 in LE Brus. Nouv. J. de Chem. 11:23, 1987.)


The increase in bandgap, �E, of a semiconductor due to size quantization isthen given by

�E � �2

�

�

2

R

2

2� � �1.

�7R9e2

� (10.2)

where the first term on the right-hand side is the localization energy [the particle-in-a-box energy of the charges in a box (more correctly, in this case, a sphere) ofradius R, modified by the reduced effective mass term) and the second term rep-resents a reduction in the energy increase due to coulombic interaction betweenthe electron and the hole and is a function of the dielectric constant of the semi-conductor, �. The increase in bandgap is inversely dependent on both the reducedeffective mass and on the square of the radius. The bandgap should therefore in-crease as a parabolic function of the decrease in size. In practice, the rate of in-crease is less than this, and an exponent considerably lower than 2 gives a betterfit of the bandgap increase with decrease in crystal size.

An obvious importance of this size quantization is that a single semicon-ductor can possess a range (sometimes a wide range) of bandgaps, which can becontrolled if the semiconductor crystal size is controlled. This “bandgap tailoring”allows (ideally) control of all properties that depend on the bandgap. The most ob-vious is the optical transmission (absorption) spectrum, and this is the propertymost often measured in quantifying size quantization, since the bandgap can be es-timated from this spectrum.

In this chapter, size quantization effects in CD films are described. Since themajority of reports on size quantization in CD films mention the effect but do notgo into detail on this aspect, as with many other chapters in this book, it will bemore efficient to tabulate the relevant literature and to deal with individual stud-ies that provide additional results of interest outside of what is included in the tableor require further discussion. CdSe and PbSe will be dealt with in a more inte-grated manner, since films of these materials, in particular CdSe, have been themost intensively studied from the viewpoint of their nanocrystallinity and quan-tum size effects.

Some words of caution in interpreting optical transmission (absorption)spectra (see also Sec. 1.4.2). Since the energy structure of nanocrystals in the quan-tum size domain is more like an atomic structure, with separate levels, than the bandstructure of bulk semiconductors, the derivation of bandgap from the absorptionspectra using the (�h�)n vs. h� plot, which is based on the density of states (bandstructure) of the bulk semiconductor, is not really valid. Nonetheless, it is oftenused to give an approximation. Since there is always a distribution of crystal sizes(and therefore of bandgaps) in these size-quantized films, there is in any case nosingle “correct” bandgap, and any value measured will be an approximation.

Another area where nanocrystallinity is important is that of surface effectsdue to the high real surface areas of such films (often tens of percent of all atoms


are located at a surface). Since these films are invariably porous to a greater orlesser extent, much of this surface is accessible to modification and treatment byliquids and gases. This aspect has been less dealt with for CD films than quantumsize effects, but some examples do exist and are briefly discussed in this chapter,with references to other sections of this book where they are of specific relevance.

The factors that influence crystal size in CD films are discussed in Section10.2 on CdSe, since they have been most studied for that material. However, theprinciples involved are general for all semiconductors.

10.2 CdSe

10.2.1 Historical Background

This chapter is the one closest to the interests of the author. These interests, bothin chemical deposition and in size quantization, were a direct result of a singleserendipitous observation. As a historical and personal interlude, this “experi-ment,” involving CD CdSe, will be briefly described and its rationale explained(in reverse order).

Accepting an invitation to spend some time in Campinas, Brazil, I plannedto utilize my experience in electrodeposition of CdSe films. Shortly before thistrip, a paper appeared describing a method of electrodeposition of CdSe, based onthe use of selenosulphate instead of the commonly used SeO2 (Ref. 4, which itselfwas based on an earlier study [5]). The trip to Campinas and the resulting abilityto spend a lot of time in the lab seemed a good opportunity to try this method. Themethod proved to be simple to reproduce. More relevantly, however, in keepingwith my overall philosophy on life, instead of clearing away the “finished” ex-periment, the beaker containing the electrodeposition solution was left sitting overthe ensuing weekend on the laboratory table to be taken care of at a more suitable(i.e., later) time. On returning to the lab after the weekend, I found the beaker inquestion to contain a bright red precipitate and also a similarly colored, transpar-ent film on the inside of the glass walls. My first thought was that this was ele-mental Se (which is usually bright red in the freshly precipitated state), based onmy understanding that selenosulphate is not very stable and can readily form ele-mental Se under certain conditions. A simple chemical test (treatment with Na2Ssolution, in which Se dissolves) showed that the precipitate was not Se. To makea fairly long story short, the next two ideas (occurring more or less at the sametime if my memory does not fail me) were either that I had discovered a new amor-phous form of CdSe or that the CdSe was size quantized, and some further exper-iments were all that were needed to verify the latter hypothesis.

Had I been more familiar with chemical deposition at that time, I might wellhave ignored the “red” CdSe as something obvious; an earlier study on CD CdSehad noted that the as-deposited films were red [6].


Just to complete this history, it should be mentioned that the weekend inquestion occurred shortly after I had spent several days in Rio de Janeiro duringCarnival. I do not make any (overt) claims that this experience affected my abil-ity to interpret the experiment in any way.

The results of this experiment and the subsequent investigation were pub-lished in Ref. 7. At the time, it was not obvious that quantum size effects wouldbe seen in strongly aggregated nanocrystals. Subsequently, this was found to oc-cur quite commonly (as seen in this chapter). While some degree of electronic iso-lation between crystals is needed for size quantization to be exhibited, this isola-tion need by no means be absolute. CdSe is a particularly attractive material toshow the phenomenon of size quantization, since its color can vary from very deepred (even black in powder form) to yellow (almost white for very tiny crystals).To give a feeling for the size dependence of the color, crystals of 6 nm will be red,4 nm orange, and �3 nm yellow, a wide range in color for a small range in crys-tal size. This variation in color can be seen from the transmission spectra shownin Figure 2.9. The spectral changes show how the absorption onset (equal to thestart of the transmission decrease) moves to the red as the deposition temperatureincreases and also as the mechanism changes from a cluster mechanism to an ion-by-ion one (see later). The bandgap range, from ca. 2.3 eV to 1.8 eV, parallels achange in crystal size from ca. 3 nm to over 12 nm (the rightmost spectrum is fora film with crystal size ca. 20 nm, but this spectrum is reached by the time the crys-tal size is close to 12 nm, where it can be considered to be bulk from the point ofview of size quantization).

10.2.2 What Determines the Size of Nanocrystalsin CD Films?

The factors that determine crystal size, a topic of particular relevance to this chap-ter, have been discussed to some extent in Section 3.4. There are two main factorsthat generally affect crystal size for any particular material: the deposition mech-anism and the deposition temperature. The hydroxide cluster mechanism is ex-pected to give a crystal size similar to that of the original hydroxide cluster (withsome growth possible as deposition proceeds). That size depends mainly on tem-perature, both because higher temperatures allow more grain growth and, possi-bly more important, lower temperatures kinetically stabilize very small nuclei insolution that are thermodynamically unstable. For example, in the hydroxide clus-ter mechanism, where crystal size is believed to be controlled mainly by the sizeof the Cd(OH)2 colloids, the relevant equilibria are

Cd2� � 2OH�D Cd(OH)2 (10.3)

A number n of these molecules can form a cluster, [Cd(OH)2]n. This cluster can


continue to grow by collecting a variety of solution species, one possibility being

[Cd(OH)2]n � 2OH�D [Cd(OH)2]n.2OH� (10.4)

[Cd(OH)2]n.2OH� � Cd2�D [Cd(OH)2]n�1 (10.5)

This may continue until eventually the cluster is large enough to be thermody-namically stable (i.e., will not redissolve). However, if the cluster is smaller thanthe critical nucleus size, then there is the possibility that the nucleus will redis-solve. The lifetime of the nucleus will then depend on its size and also on the tem-perature; lower temperatures will slow the redissolution step. Thus lower temper-ature increases the chance that a subcritical nucleus will eventually grow to astable size rather than redissolve. This kinetic stabilization of small nuclei resultsin a greater total density of nuclei and therefore smaller crystal size, since the to-tal quantities of reactants are fixed.

For the ion-by-ion reaction, nucleation is generally slower and the densityof nuclei smaller. Additionally, growth occurs (ideally) only at a solid surface;therefore nucleation is confined to two dimensions, in contrast to three dimensionsfor the cluster mechanism. The crystal growth may terminate when adjacent crys-tals touch each other or by some other termination mechanism, e.g., adsorption ofa surface-active species. These factors should be valid regardless of whether themechanism proceeds via free chalcogenide ions or by a complex-decompositionmechanism.

The effect of temperature and mechanism on the optical spectra, through thecrystal size, is clearly seen in Figure 2.9. In particular, the difference in crystal sizebetween the two rightmost spectra, both deposited at 80°C but one through thecluster mechanism and the other (HC) through the ion-by-ion mechanism, is rela-tively large: 8.5 nm for the film deposited by the cluster mechanism and 20 nm forthat deposited by the ion-by-ion mechanism. While the effect of temperature isgradual, that of mechanism is sudden. It is determined by the conditions that sep-arate the formation of metal hydroxide colloids from a solution with no metal hy-droxide phase. Decrease in the complex:metal concentration ratio and increase intemperature and pH will all favor hydroxide formation. This sudden transition onvarying the complex:metal ratio is shown, for two different solution temperatures,in Figure 10.2. The spectra are independent of the NTA:Cd ratio (NTA, nitrilotri-acetate, the complex used) until the transition between hydroxide-containing andhydroxide-free solutions is reached, whereupon they suddenly undergo a red shift(increase in crystal size, decrease in bandgap) and then no further change as theNTA:Cd ratio increases further. The crystal sizes in these films were also shown,by XRD, to change only at the NTA:Cd ratio where the spectrum changes and the experimental results agreed with thermodynamic calculations on the region of


existence of Cd(OH)2 as a function of temperature, pH, and solution composition[8].

Illumination, by light that is absorbed by the growing semiconductor crys-tals, has been shown to increase the crystal size somewhat, seen as a red shift inthe optical spectrum and decrease in bandgap by as much as 0.2 eV [9–11] (seeFig 4.3). This is probably due to photoelectrochemical reactions taking place atthe crystal surface. The chemical deposition solution can also be used to elec-trodeposit CdSe. Electrons (either from an external source, as in the case of elec-

FIG. 10.2 Optical transmission spectra of CdSe films deposited at 10°C (top) and 80°C(bottom) with varying NTA:Cd molar ratios (shown in the figures). The bandgaps esti-mated from the spectra are indicated by the thin vertical lines. (Adapted from Ref. 8.)


trodeposition, or generated by illumination during CD) can reduce selenosulphateto selenide, which reacts with Cd ions, causing (photo)electrodeposited CdSe toform on the crystals and thereby increasing the crystal size. Only light that is ab-sorbed by the semiconductor can cause this effect, as expected based on the pho-toelectrochemical mechanism. In addition, a certain minimum intensity is neededto cause a measurable red shift, but the effect saturates at high light intensities. Atlow intensities, recombination was assumed to remove the photogenerated elec-tron before it had time to reduce selenosulphate, hence the threshold. As intensityincreases, there is an increasing likelihood of photoelectrochemical reduction. Thesaturation was explained by assuming that only one electron/hole pair was effec-tive; further increase in intensity has no further effect [10,11]

Since size-quantized CdSe can undergo visible color changes with changein crystal size, this effect of illumination can clearly be used to form patterns onthe film by illuminating the film through a mask. We have even observed inter-ference fringes at the edges of these patterns, alternating between orange (nonil-luminated) and brown-red (illuminated), corresponding to destructive and con-structive interference, respectively, at the pattern edges.

It is a general observation when quantum size effects are observed in CDfilms that the blue shift is reduced somewhat as the film thickness is increased, andthis has been shown clearly for CdSe [9,11–13]. From absorption spectra, a dif-ference of ca. 0.08 eV (a crystal size difference of 0.26 nm for a crystal size of ca.4 nm) was shown to occur between thin and thick films (growth time between 10and 190 hr) [11]. In another study, using the same basic deposition solution(NTA/selenosulphate), red shifts in the photoluminescence spectra could be cor-related with a change in crystal size from ca. 4.5 nm (2.1 eV) to ca. 8 nm (1.85 eV)[12].* The widths of the photoluminescence peaks increased as the depositiontime (and therefore crystal size) increased. This was explained as an increase insize distribution of the crystals as deposition proceeded. Since the change inbandgap with crystal size increases as crystal size decreases, wider peaks in thephotoluminescence spectra, if due to increase in size distribution, means that thesize distribution increases greatly with increasing deposition time. If no change inthe size distribution were to occur, the peak widths should actually decrease withincrease in average crystal size. While photoluminescence peak widths of size-quantized films may certainly be influenced by size distribution, other factors, inparticular recombination from surface sites with various spatial separation, willalso affect the width. A microscopic study of such films would be the most de-pendable way to measure size distribution.

* The values of crystal size reported in this study varied between 6.6 and 10.6 nm. These values werebased on the effective mass approximation, which overestimates the bandgap increase with decreasingcrystal size, particularly for small sizes.


Increase in size distribution with increasing film thickness (or depositiontime) is expected for a number of reasons. One is the obvious one that, since filmgrowth can involve both crystal growth and new nucleation, the chances of anyparticular crystal growing increases with time (film thickness), due to depositionof new material on the crystal. Also, as deposition proceeds, the solution compo-sition changes, and this can lead to changes in crystal size. There are two maincauses of this. One is that, while the Cd concentration decreases, that of the com-plex does not change (this may not be the case where a volatile complexant, likeammonia, is used in an open bath). Therefore the complex:Cd ratio increases dur-ing deposition, and at some point the mechanism may change from hydroxidecluster to ion-by-ion; the latter normally gives larger crystals and may also occuron already existing crystals. The second cause is that, even assuming no change indeposition mechanism, crystal size grows slightly with decrease in reactant con-centrations; reducing the concentration of all reactants to one-half the originalconcentration resulted in a bandgap increase of between 0.05 and 0.10 eV, corre-sponding to a size increase of ca. 10% [9]. This can be explained based on the con-cept of kinetic stabilization of small nucleii. As described earlier, small nucleii arestabilized at lower temperatures, thus providing a greater chance of growth to asize where the crystal will be thermodynamically stable and resulting in a smallercrystal size. However, the lower the concentration of reactants (selenosulphateand free Cd2�, the latter of which will decrease due to both decrease in total Cdconcentration and increase in complex:Cd ratio), the slower will be this growthand therefore, as with increased temperature, the greater likelihood of redissolu-tion of the small nuclei, resulting in a larger final crystal size. Of course, this con-centration effect is important not only in the context of varying film thickness (de-position time) but also as a means for further control of crystal size.

Thicker films with crystal size more characteristic of thin films could beformed by depositing on an existing thin film from a new solution and repeatingthis to the desired thickness. This suggests that the main reason for increase incrystal size with continuing deposition is not simply because of deposition on al-ready-deposited crystals, but because of changes in the composition of the depo-sition solution. This is also borne out by a comparison of modulated electrotrans-mission (ET) and electroreflection (ER) spectroscopies [where modulation of thepotential of the film in an electrolyte results in corresponding changes in the ab-sorption and reflection) of a CD CdSe film (Fig. 10.3)]. The bandgap of the filmis ca. 2.1 eV, but the ET spectrum is broader and shifted mainly to the high-energyside compared to the ER spectrum. Since the ER spectrum samples the near-sur-face region of the film and the ET spectrum the total film, this difference suggeststhat the film is composed of smaller crystals close to the substrate and larger onestoward the surface.

The effect of film thickness was very evident for films deposited from bathsbased on N,N-dimethylselenourea and complexed with both ammonia and either


citrate or tartrate [14]. Bandgaps estimated from the transmission spectra variedby at least 0.2 eV over a thickness range from under 100 nm to a few hundrednanometers. (Similar thickness effects were seen using photoluminescence in sim-ilarly prepared films [13].) Films from the tartrate bath were deposited morerapidly than those from the citrate bath [15], and the higher bandgap found for thecitrate-based films in this study was probably due to the fact that those films werethinner than the tartrate ones. For comparable film thickness, the bandgaps weresimilar (derived from a comparison of deposition rate [15] and spectra as a func-tion of deposition time [14]). No XRD patterns were found for these films; there-fore crystal size was not directly measured. It was noted that the deposition solu-tion color changed during deposition from colorless through turbid yellow,orange, to orange-red. This color change is typical for low-temperature, hydrox-ide-cluster-mechanism CD of CdSe in general when carried out at relatively lowtemperatures.

Annealed films deposited from a N,N-dimethylselenourea/citrate/ammoniabath were shown to exhibit a (0001) XRD reflection at 2� � 13°, a reflection nor-mally forbidden in hexagonal CdSe [the (0002) reflection is the one normallyseen] [13]. This was explained by a breaking of the selection rules due to the smallcrystal size. Interestingly, this peak was very weak in thin films and prominent in

FIG. 10.3 Modulated electrotransmission (ET) and electroreflectance (ER) spectra of aCD CdSe film deposited from an NTA/selenosulphate bath at 30°C. The experiments werecarried out in an electrolyte containing sulphide and sulphite (the latter to prevent forma-tion of colored polysulphide) at a pH of 10 (buffered with NaH2PO4).


thicker ones. A broad Raman band at 250 cm�1, which is not observed normallyin CdSe and disappeared after annealing, was also observed in these films and at-tributed to a surface optical mode.

10.2.3 Photoluminescence (See also Sec. 4.2.7.1)

Photoluminescence was mentioned earlier in connection with studies in change ofcrystal size with film thickness. On a more general note, a number of photolumi-nescence studies have been reported on size-quantized CD CdSe films. Spectrashowing both bandgap luminescence (this is probably not true bandgap but re-combination from shallow surface states) [7,12,16] and dominant deep-trap lumi-nescence [10,11,17,18] have been reported. One study, using dimethylselenoureainstead of the more common selenosulphate, found both “bandgap” luminescence(which red-shifted on increasing film thickness, explained by increasing crystalsize) and a lower-energy peak at ca. 1.75 eV, attributed to larger, weakly quantum-confined crystals. It is very possible that this low-energy peak arises from surfacestates, since there was no evidence of a bimodal size distribution that would leadto two separate peaks

The role of water adsorbed on the surface of these nanocrystals in passivat-ing surface states was discussed in Section 9.2.2.2. This is seen in luminescence bychange in the spectrum from deep-trap-dominated in a dry ambient to near-bandgap-dominated in a humid ambient [17]. A study of the deep-trap lumines-cence showed that the luminescence originated from recombination of trappedcharges [17] (see also surface photovoltage spectroscopy measurements in Ref.18). From a consideration of the optically detected magnetic resonance (ODMR)signals, it was shown that the recombination occurred from sites of low symmetry,i.e., at the surface of the nanocrystals [19]. However, based on time-resolved pho-toluminescence and transient absorption measurements, even the near-bandgap lu-minescence was believed to result from shallow-trapped carriers. The dynamics ofthe photogenerated charge trapping for these films showed, using transient ab-sorption measurements, fast (subpicosecond) electron trapping to shallow surfacestates [10,20] and slower, but still relatively fast (ca. 50 ps) emptying of these shal-low traps, either to the ground state or, more likely, to deeper traps [20]. Even tak-ing into account this strong effect of water adsorption, in the author’s experiencethe luminescence from these CdSe films can be highly variable, not only in whetherit is dominated by near-bandgap or deep-trap recombination but, maybe even moreso, in the intensity of the luminescence, varying from relatively strong lumines-cence (visible to the eye) to no measurable signal at all. While the reason for thislarge variation is unknown, it is reasonable to assume that it is related to the natureof the surface of the individual nanocrystals.

A collection of studies on size quantization in CD CdSe films, together withsome relevant data, is presented in Table 10.1.


10.2.4 Annealed Films

It appears that most reported CD CdSe films are size-quantized, with crystal sizesof �10 nm. There are some exceptions. Films deposited via an ion-by-ion mech-anism at high temperature possess larger crystal size and show no size effects [8].Films deposited from an ammonia/selenosulphate bath were reported with abandgap, measured from the absorption spectrum, typical of bulk CdSe (ca. 1.74eV) [22]. The CdSe in this study was grown at 80°C from a solution containing 48mM Cd and 2.1 M NH3. Taking into account the additional complexing power ofthe selenosulphate, such a solution may be close to the transition between a hy-droxide mechanism and an ion-by-ion one.

One technique commonly used to illustrate size quantization in these filmsis annealing them to increase crystal size. This results in a gradual red shift in thespectra until eventually “bulk” CdSe (ca. 11 nm in size) is reached, after which nofurther shift, at least not one due to size, is seen (see Refs. 7, 13, 14, 15, and 23).The increase in crystal size with annealing depends on temperature and time of an-

TABLE 10.1 Possible Quantum Size Effects in CD CdSe

Max. Eg Min. crystal(eV) size (nm) Miscellaneous Refs.

2.3 ca. 3.5 NTA/selenosulphate 7, 8, 9, 17, 19�2.3 Citrate or tartrate/ammonia/ 14

DMSea

2.1 Triethanolamine/ammonia/ 21selenosulphate

2.28 Citrate/ammonia/DMSe 152.06 Tartrate/ammonia/DMSe 152.06 Citrate/selenosulphate 15

ca. 2.1 ca. 4.5 (6.5)b NTA/selenosulphate 12ca. 2.05c 6.4d NTA/selenosulphate 20ca. 2.3 Citrate/ammonia/DMSe 13

2.3 4.25 NTA/selenosulphate 112.28 4.34 NTA/selenosulphate 10

Bulk bandgap � 1.73 eV (wurtzite), ca. 1.8 eV (zincblende). All bandgap values given in this chapterare room-temperature values.a N,N-Dimethylselenourea.b This reported size is probably overestimated; simple effective mass approximation used in its esti-mation.c The optical absorption spectrum consists of a reasonably sharp onset corresponding to a bandgap ofca. 2.05 eV and a broad tail to longer wavelengths (onset of between 1.8 and 1.9 eV).d The films in this study were thick (ca 1 �m), and therefore the average crystal size may be expectedto be larger than most films, which are usually in the range of 100–200 nm thick.


nealing as well as on the material annealed and on the annealing atmosphere. ForCdSe and CdS, which are usually annealed in air or sometimes in an inert atmo-sphere, as a rule of thumb, increase in crystal size is slow up to a temperature ofca. 300°C and increases greatly at a temperature somewhere between 300 and400°C (with further growth at higher temperatures), together with a phase trans-formation (if the original film is sphalerite) to wurtzite structure. This can be seenin Figure 10.4, which shows a CdSe CD film as deposited and after sequential an-nealing treatments. (See also Fig. 1a in Ref. 24, which shows XRD spectra of es-sentially the same process). Although the size increase at low annealing tempera-tures is small, the red spectral shift occurs mainly in this region, since only a smallincrease in crystal size is necessary to obtain an appreciable red shift.

10.3 CdS

Films described in Table 10.2 are assumed to be deposited using a standard bath(Cd2�/ammonia/thiourea/relatively high temperature) unless described otherwise.

FIG. 10.4 TEM micrographs of a nanocrystalline CdSe film as deposited (upper left)and after air-annealing for 20 mn at increasing temperatures up to 500°C (lower right). The50-nm scale is the same for all micrographs except for the 450°C and 500°C ones, whichare marked by the 100-nm scale. (From S. Gorer and G. Hodes, unpublished results.)


The first report of size quantization in a CD semiconductor film was in 1981for very thin films of CdS deposited from CdSO4 and thioacetamide [25]. Quan-tum shifts were measured by photoluminescence (absorption spectra in thebandgap region could not be measured, probably because the films were ex-tremely thin). Compared to an exciton luminescence peak at ca. 507 nm (2.45 eV)for bulk material, a peak at 468 nm (2.65 eV) was obtained for ultrathin films (par-ticles of ca. 10 nm separated from each other by ca. 20 nm) and between 470 and495 nm (2.64–2.51 eV) for thicker films (thickness not defined). The peak of theultrathin films did not change if the film was heated at 800°C in H2S, a treatmentthat would almost certainly result in crystal growth well beyond the size quanti-zation limit for thicker films. The 10-nm particle size is at least twice the sizeneeded to see the observed blue shift in the spectra, suggesting that either the par-ticles were composed of a few aggregated crystals or the vertical dimensions ofthe particles were much smaller than the measured lateral ones.

Blue shifts in very thin films (�3 nm average thickness) were also more re-cently measured for CD CdS films deposited in an ultrasonic bath [34].

In general, very thin films would be expected to exhibit size quantization ifthe film thickness is less than the Bohr diameter, since quantization will occur inat least one dimension. However, in most cases, such very thin CD films can be

TABLE 10.2 Possible Quantum Size Effects in CD CdS


2.66 From thioacetamide bath, 25ultrathin layer

2.63 CdI2 262.5 Eg larger with increase in 27

thiourea concentration2.64 See text Citrate/ammonia bath 28

ca. 5 NTA bath; cluster mechanism 8ca. 2.7; see text 4.8 Standard bath 29

2.60 �20 nm 65–85°C 302.5 Cd acetate 312.69 4.1 Electrochemical-induced and 32

thiol capping agent2.58 Citrate/ammonia bath 333.2 Ultrathin layer (2–3 nm), 34

deposited in ultrasonic bath5.0 — 35

Bulk bandgap � 2.43 eV.


expected to be composed of islands, as occurs in the films described earlier, andsize quantization is not then expected if the individual crystal size is greater thanthe Bohr diameter.

In Ref. 26, optical transmission spectra of films deposited using CdI2 in aNH3/NH4

�/thiourea bath were substantially blue-shifted (to 2.63 eV) from the nor-mal CdS absorption. While no crystal size was given, from the XRD spectrumgiven in the paper it appears that the crystal size was considerably larger than 10nm, and therefore the cause of the blue shift is not clear. No such shift was seenfor films deposited under similar conditions but using CdCl2 instead of the iodide.

In Ref. 27, small increases in bandgap were found when high thiourea con-centrations were used. This was explained, in general terms, as a decrease in grainsize with increase in deposition rate.

In Ref. 28, using a citrate/ammonia bath, clear increases in bandgap (up to2.64 eV) were observed. A crystal size of ca. 4 nm was calculated from the XRDdata, and this size would be consistent with the observed optical spectra. Most ofthe XRD peaks, while riding on a strong and noisy background, do appear to bemuch sharper than would be expected for a 4-nm size. On the other hand, onehigher-angle peak [(11.0)] is considerably wider, although it remains wide evenafter annealing at 450°C for one hour, a treatment that normally will cause crystalgrowth into the size range of hundreds of nanometers. The main differences be-tween this bath and many standard baths are the use of citrate together with am-monia (citrate is a weaker complexant for Cd than is ammonia, although it mightact as a surface blocking agent through adsorbed carboxylate groups) and the highthiourea:Cd ratio (ca. 17).

In Ref. 8, crystals ca. 5 nm in size were deposited from a nitrilotriacetate(NTA)-complexed bath (no ammonia) at 40°C (a lower temperature than mostCdS depositions). The composition of the bath was such that Cd(OH)2 was pre-sent as a colloidal phase (cluster mechanism–see Chap. 3). Under conditionswhere no hydroxide phase was present and the reaction proceeded via an ion-by-ion mechanism, much larger crystals (�70 nm) and a red-shifted spectrum werefound. See Section 10.2.2 for more detail on the dependence of crystal size on thedeposition mechanism.

In Ref. 29, bandgaps up to 3 eV were reported. From the optical spectra, itis more probable that the maximum bandgap value is ca. 2.7 eV. The bandgap de-creased strongly with increase in film thickness up to 250 nm, at which stage thebandgap was close to the bulk value. While most films were deposited at a pH of11.7, it was noted that smaller particle size was obtained at pH�12 and larger onesat pH � 9.

In Ref. 30, small increases in bandgap (up to a bandgap of 2.6 eV) were ob-tained for CdS deposited in the presence of a magnetic field. The main factor de-termining the higher bandgap was a lower deposition temperature (65°C). How-ever, since the minimum crystal size, measured by XRD, was at least 20 nm, and


bandgap increase due to size quantization in CdS will only become apparent wellbelow 10 nm, there is probably another reason for the apparent increase inbandgap.

In Ref. 31, using CdAc2, a blue shift of nearly 0.1 eV (to 2.50 eV) was ob-tained, compared to films using CdCl2. In contrast to the CdCl2 deposition, whereXRD showed sharp peaks, no pattern was observed for the acetate films (amor-phous or small crystal size).

In Ref. 32, using a process of electrochemically induced CD of CdS (seeSec. 4.1.6.6), nanocrystalline CdS films were deposited by using 2-mercap-toethanol as a strongly adsorbing growth-termination (capping) agent. By in-creasing the concentration of mercaptoethanol, crystal size was reduced to the re-gion where quantum size effects were observed. Crystal sizes down to 4.1 nm(bandgaps up to 2.69 eV) were obtained. Above ca. 10 mM mercaptoethanol con-centration, film formation was prevented. It is particularly interesting (and fortu-nate) that, up to this concentration, film thickness was not seriously affected bythe mercaptoethanol. A strongly adsorbed surface layer on the crystals might beexpected to prevent adhesion between crystals and between crystals and substrate,necessary to form a film (and apparently does so at high enough concentration).This technique of surface capping by strong adsorbents, well known for semicon-ductor colloids, can therefore be applied—albeit with greater limitations expectedon the maximum adsorbent concentrations—to control particle size in CDnanocrystalline films.

In Ref. 33, a bandgap of 2.58 eV was obtained from a citrate/ammonia bath(deposition solution similar to that used in Ref. 28, only with ten times the Cd con-centration). No crystal size measurement was given.

A transient optical absorption study of charge recombination dynamics wasrecently reported on CD CdS with a crystal size of 5 nm (no details of the deposi-tion were provided) [35]. Three different time constants for the decay of the transient absorption bleaching were measured (0.8 psec, 17 psec, and 800 psec)and were attributed to three groups of nanocrystals with specific defects. Tri-octylphosphine oxide (TOPO), a well-known passivating agent for CdSenanocrystals that acts by binding to Cd, was found to increase the relative contri-bution of the 800-psec recombination, which was therefore attributed to volumerecombination, since the TOPO is expected to reduce the relative contribution ofsurface recombination (represented by the two shorter time constants).

10.4 PbSe

The other semiconductor, apart from CdSe, that has been studied with deliberateemphasis on quantum size effects, is PbSe (see Table 10.3). There are several rea-sons for this. One is the long-known use of CD to deposit PbS and later PbSe. Sec-ond, and of particular importance, the electron/hole effective mass in PbSe is very


low (0.034 times the free-electron mass at room temperature); from Eq. (10.2),this means that the quantum size effect will be particularly large, i.e., the bandgapwill increase noticeably even for relatively large crystals and the increase for smallcrystals will be very large. Third, in keeping with the usual small crystal size ob-tained by CD, TEM images of CD PbSe from an earlier study showed crystals ofthe order of 10 nm in size [36]. Motivated by these considerations, a study of CDPbSe films was carried out with an emphasis on crystal size, the control of thissize, and size quantization effects in the films [9,37–39].

Much of this study parallels that of CdSe; the reader is referred to Section5.3.2 for a discussion of these films. Summing up, crystal size can be controlledby deposition temperature (lower temperature gives a smaller crystal size) and bythe dominant deposition mechanism (ion-by-ion results in larger size than clustergrowth), for the same reasons as described for CdSe. Annealing can also be usedto increase crystals size, which results in red shifts of the absorption spectra [39].However, there are also some differences between PbSe and CdSe films. PbSesometimes grows with a bimodal size distribution, i.e., areas of small crystals andothers of considerably larger ones. Also, PbSe forms very large (ca. 1 �m) crys-tals under conditions of high deposition temperature and ion-by-ion growth, incontrast to CdSe, where the crystal size is typically limited to tens of nanometers.The bulk bandgap of PbSe is 0.28 eV. Values of bandgap up to 1.5 eV for a crys-tal size of ca. 4 nm have been measured [39]. Figure 10.5 shows experimental val-ues for PbSe films of bandgaps, measured from absorption spectra against crystalsize (horizontal bars that show the size distribution), as well as theoretical calcu-

TABLE 10.3 Possible Quantum Size Effects in CD PbS and PbSe


PbS1.97? 3a Pb2� � H2S 47

PbSeca. 0.67b; ca. 30 Selenourea or selenosulphate 43, 440.35c

1.5 3.5 Selenosulphate and various 9, 37, 39complexants

�0.8 Amorphous Selenosulphate � thiosulphate 46

PbSe bulk bankgap � 0.28 eV.a Size of coprecipitated PbS.b Photoconductivity maximum; see text.c From low-temperature photoluminescence spectrum and converted to room temperature.


lations of the bandgap–size relationship based on a hyperbolic band model [40]and an envelope function calculation [41]. The results agree for the most part withthe former model. Such agreements between theory and experiment for size quan-tization should be treated with some caution. Thus, more recent experiments onPbSe nanocrystals in a phosphate glass matrix suggest a smaller increase inbandgap with decreasing crystal size [42]. For example, a crystal size of 3.5 nm,which exhibited an apparent bandgap of ca. 1.5 eV in the CD films, showed oneof ca. 1.0 eV in the glass samples. Several reasons could be put forward to explainthis discrepancy. The surrounding glass matrix could reduce the confinement inthe glass samples, although it might be argued, with equal validity, that the con-tact between the nanocrystals in the film does the same thing. Additionally, thelead chalcogenides usually exhibit weak absorption near (and even quite far from)the bandgaps, and it is sometimes difficult to unambiguously determine thebandgap in a thin film.

FIG. 10.5 Experimentally measured values of bandgap of PbSe films (horizontal bars:The length gives the experimental uncertainty in size, mainly due to the size distribution).The broken curve gives the theoretical relationship between bandgap and crystal size basedon the hyperbolic band approximation used for PbS in Ref. 40. The room-temperature re-duced effective mass (0.034) was calculated from the low-temperature value (0.022) (R.Dalven, Infrared Phys. 9:141, 1969.) according to the temperature dependence given in H.Preier, Appl. Phys. 20:189, 1979. The dotted curve is a more recent calculation based on anenvelope function calculation [41].


These films were active as electrodes in a photoelectrochemical cell (see Sec.9.2.4). Figure 10.6 shows the photocurrent spectral response of four different PbSefilms with different crystal sizes (the long-wavelength cutoff at 1.4 �m is due toabsorption by water from the aqueous solution used). Spectrum (a) is for very small(�4 nm) crystals with a bandgap of ca. 1.5 eV. Spectrum (b) is for a film withmainly small crystals and also areas of larger crystals (bimodal distribution). Thelonger-wavelength feature is due to the larger crystals. Spectrum (c) represents amixture of moderately small and moderately large crystals, while (d) contains onlyfairly large crystals. The important point to be made is the control over the shapeof the spectral response (the details of the crystals sizes and preparation methodsare less important and are not dealt with here). These spectral responses can betranslated, in principle, into spectral responses for detectors based on these films(in practice, these films, used as photoelectrodes in the present setup, would prob-ably be neither efficient nor very stable as detectors). Since PbSe photodetectorspeak in the infrared and their sensitivity decreases strongly toward the visible, theability to control both the crystal size (� bandgap) and size distribution (range ofspectral response) allows extensive tailoring of the spectral response of a detectorbased on these films. Response (c) in Figure 10.6 is particularly interesting, for itindicates the possibility of a flat response over a wide range.

Blue shifts in photoconductivity spectra, attributed to quantum size effects,were noted in PbSe films deposited using either selenourea or selenosulphate

FIG. 10.6 Photoelectrochemical photocurrent spectra (in a selenosulphate electrolyte)of CD PbSe films deposited under various conditions with four different crystal sizes (seetext). The cutoff at 1400 nm is due to absorption by the water. (From S. Gorer, PhD dis-sertation, Weizmann Institute of Science, Rehovot Israel, 1996).


[43,44]. The peak in the photoconductivity spectrum shifted from ca. 1.7 �m forthe as-deposited films to 3.6 �m after annealing. These peak values do not give avalue of the bandgap directly, but an approximate value of 0.67 eV for the as-de-posited film may be estimated, assuming that the shapes near the photocurrent on-set of the as-deposited and annealed photoconductivity spectra are similar. Crys-tal size in similar films was measured to be ca. 30 nm [43,45], slightly less for theselenourea samples and slightly more for the selenosulphate ones. Only smallbandgap shifts (�0.1 eV) are expected for this crystal size. Shifts in photolumi-nescence peaks of ca. 0.02 eV (selenosulphate bath) and 0.07 eV (selenourea bath)were more in line with the measured crystal size, particularly taking into accountthat the crystal size of the former was somewhat larger than that of the latter [44].In comparing these results (using the selenosulphate bath) with those earlier,which give larger quantum effects, two differences are relevant. One is that theformer are considerably thicker (by at least several times) than the latter. This in-creases the crystal size in general and particularly so for many of the PbSe films.The second is that without further information, neither photoconductivity nor pho-toluminescence spectroscopy (in particular when peaks, rather than onsets, areused) are accurate indicators of bandgap, although they can provide indication ofquantum size effects.

When thiosulphate was added to the selenosulphate deposition bath, noXRD pattern was observed in the films, and they were presumed to be amorphous[46]. Annealing at 350°C induced crystallization. Just one minute at this temper-ature was enough to give a definite XRD pattern. The crystal size after this treat-ment was 13 nm, meaning that the size before annealing, if any crystallinity didexist, was considerably less than this. A strong blue shift in the absorption spec-trum of the as-deposited film showed an onset of ca. 0.8 eV.

Finally, two related studies can be mentioned here. It was noted that when aquartz plate was immersed overnight in a solution of Pb(ClO4)2 and poly(vinyl al-cohol) through which H2S had been bubbled, a film formed on the plate parallelwith formation of a colloidal PbS sol [47]. The film was extremely thin (maximumabsorbance of �0.015 at 400 nm). The absorption spectrum of this film was sim-ilar to that of the PbS sol and consisted of several absorption peaks with an ab-sorption onset of ca. 630 nm (1.97 eV). It is not clear that this is the true bandgaponset, for the same reasons as discussed previously (weak absorption close to thebandgap). The XRD crystal size of the precipitate was ca. 3 nm.

It is interesting that, apart from this study, quantum size effects have not beendescribed in CD PbS films, in contrast to PbSe ones. Although PbS does showweaker quantum effects than does PbSe (because of its larger effective mass), itstill should show strong quantum size effects—greater than CdSe, for example. Forsome reason, PbS seems to grow with larger crystal size than many other semi-conductors. However, there is no a priori reason to indicate that size-quantized PbScould not be deposited by CD, and it is likely that an effort to do so would bear fruit.


The second study of possible relevance reported that PbSe, precipitatedfrom selenosulphate solution (not in the form of a film), was found to have an(electrical) bandgap, measured by temperature-dependent resistivity, of 0.4 eV[48]. In the same study, samples prepared by reaction of solid lead tartrate withH2Se exhibited an electrical bandgap of 0.92 eV. These results suggest the occur-rence of size quantization.

10.5 OTHER SEMICONDUCTORS

Quantum size effects have been noted, either explicitly or implicitly (in somecases, not noted in the study but inferable from the optical spectra), in a numberof other CD semiconductor films. In most cases, the quantum size effects are aside issue and are only briefly mentioned. For that reason, the results will be givenin Tables 10.4 and 10.5, with additional information given for specific studieswhere available.

10.5.1 ZnS

In Ref. 49, the bandgaps of these ZnS films were dependent on the deposition tem-perature and varied from 4.05 eV (3°C) to 3.88 eV (90°C) (bulk cubic bandgap �3.6 eV). The crystal size measured by TEM for the 90°C deposition was 6–8 nm,in the range, if somewhat large, for the observed bandgap increase in ZnS. For anacidic bath (no ammonia and higher Zn concentration), the bandgap was the 3.6eV of bulk cubic ZnS.

No mention of quantum size effects was made in the study in Ref. 50, andin fact the measured bandgap (3.76 eV) was reported to agree with the literaturevalue. The literature value meant was probably that of the hexagonal form of ZnS,which has a somewhat higher bandgap (ca. 3.8 eV) than the cubic form. However,electron diffraction showed the cubic form was deposited. Therefore this value isindicative of size quantization.

In Ref. 51, the films were reported to contain significant amounts of hy-droxide or oxide. This is a common property of CD ZnS films deposited from al-kaline solution (see Chap. 3).

In Ref. 52, no XRD pattern was observed in the as-deposited films. An-nealing at 450°C resulted in a bandgap of 3.7 eV (and conversion of some of theZnS to ZnO). This value suggests that hexagonal ZnS is formed (at least, after an-nealing).

10.5.2 ZnSe

In Ref. 53, from the optical absorption spectrum, a slight blue shift (2.63 eV com-pared to bulk 2.58 eV) was explained by size quantization. Note that XRD of theprecipitated powder suggests a crystal size of ca. 3 nm (no pattern was discerned


in the film, possibly due to the small amount of material together with the smallcrystal size). Such a small size would be expected to result in a larger bandgap.

No reference to size quantization was made in Ref. 54. However, thebandgap is somewhat larger than for a nonquantized sample. The crystal size (ca.10 nm from both electron microscopy and X-ray diffraction) is a bit large to showsize effects in the spectra. In Table 10.4, there is some degree of inconsistency be-tween the various bandgaps and crystal sizes, if it is assumed that size quantiza-tion determines the bandgap. It may be that other factors influence the bandgaphere.

In Refs. 55 and 56, the Zn:Se atomic ratio of 3:2 found in the films, togetherwith the TEM micrographs, which show very small ZnSe crystals, suggest thatsize-quantized ZnSe is formed together with ZnO [or Zn(OH)2].

TABLE 10.4 Possible Quantum Size Effects in other CD II-VI Semiconductors.


ZnS4.05 (3°C) — Thioacetamide/ammonia 493.88 (90°C) 6–8 493.76 Thiourea/ammonia/hydrazine 503.85 (85°C) 3 Thiourea/ammonia/amines 513.9 Thioacetamide/triethanolamine/ 52

ammonia

ZnSe2.63 ca. 3 nma DMSb/citrate/ammonia/50°C 532.7 �10 Selenourea/ammonia/ 54

hydrazine/60°C2.9 2–2.5 Selenourea/ammonia/ 55,56

hydrazine/70°C

HgSca. 2.3 (0°C) 3–4 Thiosulphate 57

1.9 (85°C) 8see text 20 Thiosulphate/ammonia 58

HgSe1.42 No XRD seen Selenosulphate/Hg0- 59

formamide-NaOH2.5 7.7 Selenosulphate/ammonia/10°C 60

ZnS (bandgap 3.8 eV hexagonal and 3.6 eV cubic), ZnSe (2.58 eV), HgS (2.1 eV cinnabar; zero-gapmetacinnabar), and HgSe (�0.15 eV)a From precipitate.b N,N-dimethylselenourea.


10.5.3 HgS

In Ref. 57, since the film thickness increased with deposition temperature [from50 nm (0°C) to 180 nm (85°C)], the decrease in bandgap with increase in deposi-tion temperature may also be due, to some extent, to this factor. From the weakXRD pattern (shown for a deposition at 28°C), the HgS appears to be the cinnabarcrystal form with a bulk bandgap of 2.1 eV. However, this would not explain thebandgap of 1.9 eV measured for the 85°C deposition.

A similar deposition, only with the addition of ammonia, was carried out[58]. This is a major change (e.g., alkaline instead of acidic conditions). While ahigh bandgap (3.1 eV) was measured from the sharp drop at 400 nm in the optical

TABLE 10.5 Possible Quantum Size Effects in Miscellaneous CD Semiconductors


Sb2S3

2.48 — Thiosulphate/SbCl3 in 67CH3COOH

1.7 �20 Annealed at 170°C 672.21 — Thiosulphate/SbCl3 in acetone/ 68

10°C/pH 51.79 �20 Annealed at 250°C 68

Cu-SCuS—2.0 �13 Thiosulphate/pH 5/50°C 69CuS—1.55 11 (200°C) Thiosulphate/dimethylthiourea 70Cu1.8S—1.55 19 Annealed at 300°CCu1.96S—1.4 20 Annealed at 400°C

Cu-Se2.38(d) 1.9(i) 20 Selenosulphate/ammonia/25°C 712.20(d) 1.3(i) 29 Annealed at 300°C (Cu1.85Se)

In2S3

2.7 (50°C) ca. 6 Thioacetamide/pH 1–3.1 (HCl) 75

Ag2Se1.8 direct 9 Selenosulphate/ammonia 77

CuInSe2

1.4 (direct) — Selenosulphate/ammonia/ 79citrate/40°C

1.15 (520°C) 20 791.3 direct 14 Selenosulphate/citrate/ 80

pH �8/25°C1.02 (300°C) 23 80


transmission, this is not likely to be a bandgap transition, for several reasons. Oneis the color of the films—golden yellow (100 nm thick) to red (500 nm). This colorwas seen also as a more gradual drop in the transmission beginning at ca. 600 nm.(An even more gradual reduction in transmission occurred to longer wavelengthsand is probably related to small amounts of lower-bandgap modifications found inthe films). Second, the XRD spectrum of the films indicated a crystal size of atleast 20 nm, making it unlikely that the transition at 3.1 eV was due to size quan-tization of cinnabar with a bulk bandgap of 2.1 eV.

10.5.4 HgSe

HgSe is normally a semimetal with a negative bandgap. However, from the twostudies noted next, either the bulk bandgap of these films is considerably larger,the quantum size effect is huge, or the measured bandgaps in all cases were highertransitions rather than the primary bandgap. More studies with careful investiga-tion of the optical properties are needed here.

In Ref. 59, no XRD pattern was observed for the as-deposited films, whichwere thus assumed to be amorphous. Annealing at 200°C was enough to give anXRD pattern (of hexagonal HgSe), although no crystal size details were given.

In Ref. 60, the HgSe was probably in the form of the cubic phase: The pre-cipitated powder was cubic, while the film was highly textured (111) [which couldalso belong to the hexagonal (0002) reflection]. The very large value of bandgapfound (2.5 eV) was compared with the value of ca. 3.2 eV previously found for 2-to 3-nm HgSe colloids [61].

10.5.5 Bi2S3

There are many reports on CD Bi2S3 where larger-than-normal bulk value (1.3eV) bandgaps are obtained. The section on Bi2S3 in Chapter 6 should be read forthose interested in quantum size effects in this material. Table 6.2, on Bi2S3, givesvalues of reported bandgaps, and values of between 1.6 and 1.9 are particularlycommon. However, as pointed out in that chapter, such values should be treatedwith caution; most of the bandgap values are calculated from an extrapolation ofthe plot of (�h�)2 vs. h� to (�h�)2 � 0 (see Sec. 1.4.2), and consideration of theraw transmission spectra, where they are given, strongly suggests that the absorp-tion onsets are often considerably red-shifted from this calculated value. This isparticularly true for films deposited from acidic solutions, which have a largercrystal size than those deposited from alkaline solution (see later).

X-ray diffraction patterns, with peak widths characteristic of a crystal sizeof at least 10 nm and usually 20 nm or more, have been obtained from Bi2S3 filmsdeposited from acidic solutions. In contrast, XRD has not shown any peaks for as-deposited films from alkaline baths, but annealing at low temperatures(150–200°C) was sufficient to give fairly sharp (equivalent to at least 20-nm crys-


tal size) XRD peaks. However, the optical absorption spectra changed only little,if at all, after this annealing (see, e.g., Refs. 62 and 63). This provides further ev-idence that the optical spectra do not reflect size quantization in general.

Taking all this into account, there seems to be no merit in including Bi2S3 inthe present tables. Reported bandgap values have already been given in Table 6.2,and crystal sizes have been summed up in the previous paragraph. Instead, someremarks on specific studies will be made.

In Ref. 64, red spectral shifts were obtained upon illumination of the filmsduring deposition. It was thought that this was due to increasing film thickness asa result of the illumination. However, consideration of the raw transmission spec-tra in this study does indicate a change in bandgap with varying illumination con-ditions (intensity and particularly duration). Estimation (by this author) of thebandgap from the transmission spectra gives ca. 1.9 eV (no illumination) to 1.5 orless eV under full solar illumination during the entire deposition. Similar illumi-nation-induced crystal growth with resulting increased bandgaps has been dis-cussed for CdSe earlier in this chapter.

Reference 65 does relate in some detail to size quantization in the Bi2S3 films.The crystal size was reported to increase and the bandgap to decrease with in-creasing film thickness. The deposition was from acidic solution, and considera-tion of the XRD pattern shown indicates a crystal size of at least 20 nm. Slow-scanXRD was interpreted to give a range of crystal sizes from 5.2 nm (for a 53-nm-thickfilm) to 8 nm (220-nm-thick), which is not in agreement with the (normal scanspeed) spectra shown. Bandgaps as high as 2.22 eV were calculated from the (�h�)2

vs. h� plots, but these do not match the values expected from the optical absorptionspectra, which indicate much lower values. A re-estimation from these (�h�)2 vs.h� plots, but at lower values of absorbance (� 104 cm�1 rather than the range closerto 105 cm�1 that was used), gives better linearity and a gap of ca. 1.4 eV, fairly in-dependent of film thickness and quite close to the literature value.

10.5.6 Bi2Se3

There are very few studies on CD Bi2Se3 (see Chap. 6) and only one that providesa reasonably in-depth study of the optical properties [66]. In this work, N,N-dimethylselenourea was used as the Se source in an alkaline triethanolamine bath.The spectrum of a thin film (90 nm thick) was quite different than the spectra ofthicker ones (150 or 250 nm, which were similar). The 90-nm film exhibited aclear onset with a direct bandgap, calculated from the (�h�)2 vs. h� plot, of 1.7 eV.(The literature room-temperature value is ca. 0.12 eV.) Annealing at 200°Ccaused a red shift in the spectrum to a calculated bandgap of 1.57 eV. For thethicker films, the as-deposited direct bandgap was 1.41 eV and after annealing,1.07 eV. For these thicker films, however, there was also loss of transmission atlower energies. Both scattering and an indirect bandgap were suggested as possi-


ble causes of this. Based on the latter possibility, indirect bandgaps between 0.2and 0.35 eV were calculated for this region of the spectra. X-ray diffraction ofthese films showed some sign of crystallinity and a very rough estimation of �10nm could be made from this spectrum. Annealing at 200°C increased the crystalsize by only a few nanometers. A rough estimation of the increase in bandgap ex-pected for a 12-nm crystal size (based on a literature value of the reduced effec-tive mass of �0.06) is ca. 0.15 eV. This suggests that the apparently indirect gapmeasured for the thicker films is somewhat quantized and that the direct gap mea-surement is a higher transition. The higher value for the thin film may be a largesize quantization effect, but information on crystal size is needed.

10.5.7 Sb2S3

In Ref. 67, the increase in bandgap of the as-deposited film was attributed to amixed amorphous-polycrystalline structure (apparently no XRD pattern wasfound for the as-deposited film). The onset of absorption in the transmission spec-trum was sharp for the as-deposited film, and any polycrystalline (nonquantized)material would be expected to give some absorption at lower energies, even if theamorphous phase possessed a higher bandgap. Therefore size quantization seemsto be a more reasonable explanation for the high bandgap of the as-deposited film.

In Ref. 68, the higher bandgap of the as-deposited films (2.21 eV vs. 1.79eV for annealed films) was explicitly explained by size quantization.

10.5.8 Cu-S

Bandgap measurements for Cu sulphides and selenides are complicated by the factthat these semiconductors are normally degenerate, with high free-carrier absorp-tion in the near-infrared and possible Moss–Burstein shifts (due to saturation ofthe top of the valence band by holes) in the optical gap. It is quite possible thatvariations in bandgaps in these materials are due to differences in stoichiometry,phase, and doping rather than to any quantum size effect. Only studies where crys-tal size can be estimated and are possibly in the quantum size range are given here.

In Ref. 69, various compositions were obtained by varying the Cu:thiosul-phate ratio. Only CuS gave an XRD pattern that allowed an estimation of crystalsize; other compositions and phases (Cu1.8S, Cu1.4S, Cu2S) showed no XRD pat-terns. The (indirect) bandgaps found for these films were 2.0, 2.0, and 1.7 eV, re-spectively.

In Ref. 70, small increases in bandgap were found for Cu1.8S (1.55 eV; lit-erature value 1.2) and Cu1.96S (1.4 eV; literature value ca. 1.3).

The apparent discrepancy between the bandgaps of CuS (both indirectbandgap values) found from these two studies should be noted. The former founda considerably larger value, even though the crystal size was slightly larger andthe film was not annealed. A comparison of the two transmission spectra does not


suggest such a large difference, and the difference may be due simply to the prob-lems in measuring bandgaps of Cu-chalcogenides discussed earlier.

10.5.9 Cu-Se

Reference 71 is the only study that explicitly invokes size quantization for Cu-Sefilms. This was based mainly on the decrease in bandgap found upon annealing.In the absence of effective mass values for the various compounds, it can only besaid that the 20-nm size measured for the as-deposited films is a bit large in gen-eral to explain the observed spectral shifts, but not decisively so. It is important tonote that the electrical resistivity of the films did not change after annealing; thisimplies that the decrease in bandgap on annealing is not due to a change in theMoss–Burstein band filling, since this would be expected to be paralleled by anincrease in resistivity.

The same study also described CuSe films, which were obtained when N,N-dimethylselenourea was used instead of selenosulphate. The crystal size in thiscase was slightly larger (23 nm). The bandgap was measured to be ca. 2.13 eV.In view of the difficulty in measuring bandgaps in these materials, this is close tothe value of 2.0 eV reported for CD CuSe films with a crystal size of 40 nm, whereno quantum size effects are expected [72].

Other studies on Cu-Se where small crystal sizes were measured and there-fore are potential candidates to exhibit size quantization give crystal sizes between10 and 20 nm and bandgaps (direct) of ca. 2.2 eV for Cu2�xSe and ca. 2.8 eV forCu3Se2 [73,74]; a value of 2.37 eV was measured for CD Cu3Se2 with a 40-nmcrystal size [72].

10.5.10 In2S3

In Ref. 75, the bandgap increased gradually from 2.32 eV to 2.7 eV as the depo-sition temperature varied from 90°C to 50°C. The pH also had an effect on thecrystal size (at 70°C, Eg was 2.56 eV at a pH of 3.1 and 2.41 at a pH of either 1.3or 1.5). This could be due to a change of mechanism [In(OH)3 is very acidic andis likely to be present at the higher pH but not at the lower one). Transmissionelectron microscopy showed the films to be composed of both needles and roundparticles. At 90°C, the round particles were 8–10 nm in size and the needles 20–30� 2 nm. At 50°C, there were fewer needles and the round particles were ca. 6 nm.X-ray diffraction showed a larger particle size: 15 nm (90°C), 13 nm (70°C), and10.6 nm (50°C). The large apparent size estimated from XRD, and the fact that an-nealing, which reduced the bandgap (see later), did not narrow the XRD peaks,suggested that the XRD spectrum was weighted toward a small fraction of largercrystals. Eg decreased gradually on annealing from 2.7 eV (50°C sample) to 2.53eV (350°C anneal). The increase in Eg was ascribed to a combination of size quan-tization and excess S (measured by EDX analysis).


The original CD of In2S3, by the same method as earlier [76] noted a crys-tal size that was probably no larger than 5 nm, although the bandgap was measuredto be 2.45 eV, similar to the direct bulk gap of In2S3. A lower energy weak ab-sorption which was seen in single-crystal material was absent in the films. In ad-dition, the photoconductivity spectral peak (of a film annealed at 250°C) was blue-shifted 0.13 eV, compared to single-crystal material.

10.5.11 Ag2Se

In Ref. 77, the (direct) bandgap of 1.8 eV was compared to the bulk value of ca.1.3 eV bulk and the difference ascribed to size quantization. The absorption spec-trum shows considerable absorption to beyond 820 nm (the maximum wavelengthmeasured), and it is not clear whether this is due to scattering or to a lower-energy(possibly indirect) absorption.

10.5.12 CuInSe2

Two cases of what are probably size quantization can be found in CD CuInSe2

films. In neither case was size quantization used to explain the larger-than-normal(1.02 eV) bandgaps. It is notable that these examples appear to be the only onesshowing this phenomenon for CuInSe2 in any form. One example of an anoma-lously high bandgap—1.38 eV—was found for CuInSe2 films sputtered onto glassat 77 K and was attributed to the amorphous structure; i.e., no XRD pattern wasobserved [78]. This could also be due to size quantization. The bandgap of thesesputtered films was calculated from a plot of (�h�)n vs. h�, where n � 1 (amor-phous semiconductor). If replotted for n � 2 (direct bandgap, crystalline semi-conductor), an even higher bandgap would result.

In Ref. 79, as-deposited films exhibited a bandgap of 1.4 eV, which droppedon annealing at 110°C to 1.225 and then fell more gradually to a final value of 1.15eV at 520°C. Crystal size measured by XRD was given only for the 520°C an-nealed film (20 nm; very small for that temperature of annealing), which meansthe crystal size of the as-deposited film was probably much smaller. Taken to-gether with the following study, this seems to be a clear case of size quantization.

The main differences in the deposition procedure between the study in Ref.80 and the previous study are the somewhat lower temperature, the (almost cer-tainly) lower pH (while not given; the previous deposition, using ammonia, wasprobably carried out at a minimum pH of 10), and the lack of ammonia in this study.

REFERENCES1. P Krumholz, O Kruh. Mikroch. 17:210, 1935.2. E Jensen. Tidsskr. Kjemi. Bergvesen 19:57, 1939.3. E Jensen. Tidsskr. Kjemi. Bergvesen 19:75, 1939.


4. M Cocivera, A Darkowski, B Love. J. Electrochem. Soc. 131:2514, 1984.5. M Skyllas-Kazacos, B Miller. J. Electrochem. Soc. 127:2378, 1980.6. RA Boudreau, RD Rauh. J. Electrochem. Soc. 130:513, 1983.7. G Hodes, A Albu-Yaron, F Decker, P Motisuke. Phys. Rev. B 36:4215, 1987.8. S Gorer, G Hodes. J. Phys. Chem. 98:5338, 1994.9. G Hodes. Isr. J. Chem. 33:95, 1993.

10. P Nemec, D Mikes, J Rohovec, E Uhlirova, F Trojanek, P Maly. Mater. Sci. Eng. B69:500, 2000.

11. F Trojanek, R Cingolani, D Cannoletta, D Mikes, P Nemec, E Uhlirova, J Rohovec,P Maly. J. Cryst. Growth 209:695, 2000.

12. R Garuthara, G Levine. J. Appl. Phys. 80:401, 1996.13. BK Rai, HD Bist, RS Katiyar, MTS Nair, PK Nair, A Mannivannan. J. Appl. Phys.

82:1310, 1997.14. MTS Nair, PK Nair, H Pathirana, RA Zingaro, EA Meyers. J. Electrochem. Soc.

140:2987, 1993.15. VM Garcia, MTS Nair, PK Nair, RA Zingaro. Semicond. Sci. Technol. 11:427, 1996.16. P Maly, J Kudrna, F Trojanek, D Mikes, P Nemec, AC Maciel, JF Ryan. Appl. Phys.

Lett. 77:2352, 2000.17. E Lifshitz, I Dag, I Litvin, G Hodes, S Gorer, R Reisfeld, M Zelner, H Minti. Chem.

Phys. Lett. 188, 1998.18. L Kronik, N Bachroch-Ashkenasy, M Leibovitch, E Fefer, Y Shapira, S Gorer,

G Hodes. J. Electrochem. Soc. 145:1748, 1998.19. E Lifshitz, I Dag, I Litvin, G Hodes. J. Phys. Chem. B 102:9245, 1998.20. XC Ai, R Jin, CB Ge, JJ Wang, YH Zou, XW Zhou, XR Xiao. J. Chem. Phys.

106:3387, 1997.21. H Cachet, H Essaaidi, M Froment, G Maurin. J. Electroanal. Chem. 396:175, 1995.22. RC Kainthla, DK Pandya, KL Chopra. Sol. State Electron. 25:73, 1982.23. MTS Nair, PK Nair, RA Zingaro, EA Meyers. J. Appl. Phys. 74:1879, 1993.24. S Gorer, G Hodes, Y Sorek, R Reisfeld. Mater. Lett. 31:209, 1997.25. GC Papavassiliou. J. Solid State Chem. 40:330, 1981.26. T Nakanishi, K Ito. Sol. Energy Mater. Sol. Cells 35:171, 1994.27. ME Özsan, DR Johnson, M Sadeghi, D Sivapathasundaram, LM Peter, MJ Furlong,

G Goodlet, A Shingleton, D Lincot, B Mokili, J Vedel. In: 1st World Conf. on PV En-ergy Conversion, Waikoloa, Hawaii, 1994, p 327.

28. MTS Nair, PK Nair, RA Zingaro, EA Meyers. J. Appl. Phys. 75:1557, 1994.29. KK Nanda, SN Sarangi, S Mohanty, SN Sahu. Thin Solid Films 322:21, 1998.30. O Vigil, A Arias-Carbajal, F Cruz, G Contreras-Puente, O Zelaya-Angel. Mater. Res.

Bull. 36:521, 2001.31. M Rami, E Benamar, M Fahoume, F Chraibi, A Ennaoui. Solid State Sci. 1:179,

1999.32. K Yamaguchi, T Yoshida, N Yasufuku, T Sugiura, H Minoura. Electrochem.

67:1168, 1999.33. A Arias-Carbajal Reádigos, VM García, O Gomezdaza, J Campos, MTS Nair, PK

Nair. Semicond. Sci. Technol. 15:1022, 2000.34. AI Oliva, O Solis-Canto, R Castro-Rodriguez, P Quintana. Thin Solid Films 391:28,

2001.


35. P Nemec, P Formanek, D Mikes, I Nemec, F Trojanek, P Maly. Phys. Status Solidi(b) 224:481, 2001.

36. RC Kainthla, DK Pandya, KL Chopra. J. Electrochem. Soc. 127:277, 1980.37. G Hodes, A Albu-Yaron. In: Proc. Electrochem. Soc. 88-14:298, 1988.38. S Gorer, A Albu-Yaron, G Hodes. Chem. Mater. 7:1243, 1995.39. S Gorer, A Albu-Yaron, G Hodes. J. Phys. Chem. 99:16442, 1995.40. Y Wang, A Suna, W Mahler, R Kasowski. J. Chem. Phys. 87:7315, 1987.41. I Kang, FW Wise. J. Opt. Soc. Am. B 14:1632, 1997.42. A Lipovskii, E Kolobkova, V Petrikov, I Kang, A Olkhovets, T Krauss, M Thomas,

J Silcox, F Wise, Q Shen, S Kycia. Appl. Phys. Lett. 71:3406, 1997.43. LP Biró, A Darabont, P Fitori. Europhys. Lett. 4:691, 1987.44. D Dadarlat, RM Candea, R Turcu, LP Biro, Zasavitskii, II, MV Valeiko, AP Shotov.

Phys. Status Solidi (a) 108:637, 1988.45. RM Candea, N Dadarlat, R Turcu, E Indrea. Phys. Status Solidi (a) 90:K91, 1985.46. LP Biro, RM Candea, G Borodi, A Darabont, P Fitori, I Bratu, D Dadarlat. Thin Solid

Films 165:303, 1988.47. S Gallardo, M Gutierrez, A Henglein, E Janata. Ber. Bunsen-Ges. Phys. Chem.

93:1080, 1989.48. NI Glistenko, AA Eremina. Zh. Neorg. Khim. 5:1003, 1960.49. R Ortega Borges, D Lincot, J Vedel. In: 11th ECPV Solar Energy Conf., Montreux,

Switzerland, 1992, p 862.50. JM Doña, J Herrero. J. Electrochem. Soc. 141:205, 1994.51. B Mokili, M Froment, D Lincot. J. de Phys. IV 5:261, 1995.52. OL Arenas, MTS Nair, PK Nair. Semicond. Sci. Technol. 12:1323, 1997.53. CA Estrada, PK Nair, MTS Nair, RA Zingaro, EA Meyers. J. Electrochem. Soc.

141:802, 1994.54. JM Doña, J Herrero. J. Electrochem. Soc. 142:764, 1995.55. CD Lokhande, PS Patil, H Tributsch, A Ennaoui. Sol. Energy Mats. Sol. Cells

55:379, 1998.56. CD Lokhande, PS Patil, A Ennaoui, H Tributsch. Appl. Surf. Sci. 123:294, 1998.57. SS Kale, CD Lokhande. Mater. Chem. Phys. 59:242, 1999.58. M Najdoski, I Grozdanov, SK Dey, B Siracevska. J. Mater. Chem. 8:2213, 1998.59. P Pramanik, S Bhattacharya. Mat. Res. Bull. 24:945, 1989.60. B Pejova, M Najdoski, I Grozdanov, SK Dey. J. Mater. Chem. 11:2889, 1999.61. JM Nedeljkovic, MT Nenadovic, OI Micic, AJ Nozik. J. Phys. Chem. 90:12, 1986.62. L Huang, PK Nair, MTS Nair, RA Zingaro, EA Meyers. Thin Solid Films 268:49,

1995.63. PK Nair, J Campos, A Sanchez, L Banos, MTS Nair. Semicond. Sci. Technol. 6:393,

1991.64. PK Nair, MTS Nair, O Gomezdaza, RA Zingaro. J. Electrochem. Soc. 140:1085,

1993.65. CD Lokhande, AU Ubale, PS Patil. Thin Solid Films 302:1, 1997.66. VM García, MTS Nair, PK Nair, RA Zingaro. Semicond. Sci. Technol. 12:645, 1997.67. I Grozdanov, M Ristov, G Sinadinovski, M Mitreski. J. Noncryst. Solids 175:77, 1994.68. MTS Nair, Y Pena, J Campos, VM García, PK Nair. J. Electrochem. Soc. 145:2113,

1998.


69. I Grozdanov, M Najdoski. J. Solid State Chem. 114:469, 1995.70. MTS Nair, L Guerrero, PK Nair. Semicond. Sci. Technol. 13:1164, 1998.71. VM Garcia, PK Nair, MTS Nair. J. Cryst. Growth 203:113, 1999.72. B Pejova, I Grozdanov. J. Solid State Chem. 158:49, 2001.73. M Lakshmi, K Bindu, S Bini, KP Vijayakumar, CS Kartha, T Abe, Y Kashiwaba.

Thin Solid Films 370:89, 2000.74. M Lakshmi, K Bindu, S Bini, KP Vijayakumar, CS Kartha, T Abe, Y Kashiwaba.

Thin Solid Films 386:127, 2001.75. T Yoshida, K Yamaguchi, H Toyoda, K Akao, T Sugiura, H Minoura. Proc. Elec-

trochem. Soc. 97-20:37, 1997.76. GA Kitaev, VI Dvoinin, AV Ust’yantseva, MN Belyaeva, LG Skornyakov. Inorg.

Mater. 12:1448, 1976.77. B Pejova, M Najdoski, I Grozdanov, SK Dey. Mater. Lett. 43:269, 2000.78. H Neumann, B Perlt, NAK Abdul-Hussein, RD Tomlinson, AE Hill. Solid State

Commun. 42:855, 1982.79. JC Garg, RP Sharma, KC Sharma. Thin Solid Films 164:269, 1988.80. PK Vidyadharan Pillai, KP Vijayakumar, PS Mukherjee. J. Mater. Sci. Lett. 13:1725,

1994.


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Chemical Solution Deposition of Semiconductor Films