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Chimie des Interactions Plasma-Surface (ChIPS), Faculty of Sciences

Ph.D. Thesis

Oxygen vacancy stabilized zirconia;

synthesis and properties

Mohsin Raza

Chimie des Interactions Plasma-Surface (ChIPS)


University of Mons

BE-7000 Mons, Belgium

Mohsin Raza

Printed by UMons, March 2017

Mons, Belgium


Ph.D. Thesis

Oxygen vacancy stabilized zirconia;

synthesis and properties

Mohsin Raza

Chimie des Interactions Plasma-Surface (ChIPS)

16 March 2017

Promoter and co-promoter Stéphanos Konstantinidis and Rony Snyders

Jury Members

Prof. Dr. Jochen M. Schneider (RWTH Aachen University, Germany)

Prof. Dr. Jean F. Pierson (University of Lorraine, France)

Prof. Dr. Roberto Lazzaroni (University of Mons, Belgium)

Dr. Thomas Godfroid (Materia Nova, Belgium)

Dr. Jerome Cornil (University of Mons, Belgium)


University of Mons

BE-7000 Mons, Belgium

i

Abstract

The crystal structure of any material is a decisive factor for controlling its properties1,2. In

this respect, zirconia (ZrO2) is a material which exists in three crystallographic phases: the

monoclinic phase which is stable up to 1205 °C; the tetragonal phase appears from 1205 °C

to 2377 °C; and finally the cubic phase is thermodynamically stable from 2377 °C to 2710

°C3. Because of the superior chemical stability, high hardness, high dielectric constant and

prominent optical properties, zirconia can be exploited for a broad range of applications such

as medical implants, oxygen detectors and as wear resistant or thermal barrier coatings

(TBCs). However, for pure zirconia, it is not possible to exploit most of the above-mentioned

applications as this is restricted by the change in volume of the zirconia-based components

due to the phase transformation upon heating and cooling, which ultimately leads to the

deterioration of the device components.

Therefore, for decades and whatever the zirconia is synthesized as a bulk material or as a thin

film coating, the cubic phase is stabilized by doping zirconia with cations such as Y (yttrium).

1 A. Van De Walle, A complete representation of structure-property relationships in crystals., Nat. Mater. 7

(2008) 455–458. doi:10.1038/nmat2200

2 G.B. Olson, Computational Design of Hierarchically Structured Materials, Science (80-. ). 277 (1997) 1237–

1242. doi:10.1126/science.277.5330.1237.

3 J.P. Abriata, J. Garcés, R. Versaci, The O-Zr (Oxygen-Zirconium) system, Bull. Alloy Phase Diagrams. 7

(1986) 116–124. doi:10.1007/BF02881546.

ii

Doing so, not only stabilizes the cubic phase but also leads to the generation of oxygen

vacancies in the zirconia lattice. This situation makes Yttria Stabilized Zirconia (YSZ) useful

for TBCs but also as an electrolyte with high ionic conductivity to be used in Solid Oxide

Fuel Cells (SOFCs). However, such stabilization strategies lead to the perturbation of the

periodic potential of the oxide-ion array, which results in higher energy barrier for O2- ions

during their diffusion to a vacant site as compared to intrinsic vacancy-doped oxides 4 .

Therefore, an intense research has been developed during the last fifteen years to promote

stabilization without incorporating aliovalent cations. Stabilization phenomena for

nanometer thick zirconia films have been related to the grain size, energy input during

growth, stresses in the zirconia film and oxygen vacancy or nitrogen atom incorporation. So

far, a consensus over what drives the phase formation and stabilization of cubic zirconia has

not been reached.

In the present study, according to quantum-chemistry based calculations, It is shown that the

cubic phase is the most stable phase if more than 3 at.% of oxygen vacancies are incorporated

in the ZrO2 lattice. On the other hand, carefully designed cold plasma-based reactive

magnetron deposition experiments allowed to control the amount of O vacancies incorporated

inside the zirconia lattice. X-Ray diffraction analysis of these oxygen vacancy doped zirconia

thin films are in remarkable agreement with theoretical predictions, hence emphasizing that

the incorporation of oxygen vacancies is the sole responsible mechanism for the stabilization

of the zirconia cubic phase. However, it is also observed that any deviation from the

optimized synthesis conditions leads to the change in film phase constitution.

4 J.B. Goodenough, Ceramic technology: Oxide-ion conductors by design, Nature. 404 (2000) 821–823.

doi:10.1038/35009177.

iii

In a second step, the thermal stability of the oxygen vacancy stabilized zirconia thin films is

addressed. X-Ray diffraction experiments, show that these films are stable up to 750 °C. The

ionic and electronic conductivity measurements highlight that the ionic conductivity is in the

range of YSZ however a colossal increase in the ionic conductivity is observed when the film

thickness is ≈ 10 nm i.e. 7.4 S.cm-1 at 725 °C. However, these coatings are electronically

non-conductive. These results highlight the possibility to use oxygen vacancy doped zirconia

films in SOFC devices. Finally, the photoluminescence analysis of OVSZ revealed the

presence of 2 emission peaks centered at 388 nm and 488 nm originating from the states

present in the band gap. This result validates the theoretical calculations data predicting the

appearance of energy states in the band gap upon O vacancy incorporation.

In conclusion, it is demonstrated that the incorporation of oxygen vacancies, and therefore

material defect chemistry plays a very important role in the phase formation and properties

of reactively sputtered thin films, particularly in zirconia. This observation could pave the

way to the development of new thin film growth strategies or to the synthesis of functional

thin films with new or enhanced properties.

iv

v

Acknowledgements

I am grateful to a large number of people who have contributed, helped and supported me in

this entire thesis work.

Stéphanos Konstantinidis and Rony Snyders, thank you very much for giving me the

opportunity to perform this thesis work, for your guidance at every step, encouragement, time

and support. I have learnt a lot from you in these last 4 years. I am very fortunate to have you

as my promoter. Thank you so much for all of this!

All the collaborators, I am very thankful to all of you for helping me out and performing the

requested measurements and calculations.

All the members of Chimie des Interactions Plasma-Surface (ChIPS), thank you all for

providing me such a nice working environment. Mattia, Claudia and Nikolay many thanks

for providing good company in and off the lab. Sabine and Dany thank you for taking care of

administrative stuff each and every time.

My friends, thank you all for being part of my life, for your support and encouragement.

My Siblings, thank you all for making my life wonderful. I can never forget the time which

we had together. I love you all, a lot!

My Parents, who pray for my success day and night, I could not have achieved this without

your support, love, and encouragement. I am very thankful for your every effort, you put to

give me a good life and to make me a good person. I love you very much!

vi

vii

Table of Contents Abstract i

Acknowledgements v

Table of Contents vii

1. Introduction 1

1.1. Background 1

1.2. Aim of the research work and strategy 5

References 7

2. Zirconia 9

2.1. The Crystal structure of Zirconia (ZrO2) 9

2.2. Stabilization strategies for zirconia c-phase 10

2.2.1. Stabilization of Zirconia c-phase by doping 11

2.2.2. Stabilization of high temperature phase of Zirconia without doping 13

2.3 Properties and applications of zirconia 22

2.3.1 Mechanical properties of zirconia 23

2.3.2 Thermal properties of zirconia 24

2.3.3 Ionic conduction of stabilized zirconia 24

2.3.4 Optical properties of Zirconia 27

References 28

3. Thin film growth 31

3.1. Sputtering 32

3.1.1. Glow discharge 34

3.1.2. Magnetron sputtering 38

3.1.3. Reactive magnetron sputtering 41

3.2. Thin film growth 44

3.2.1. Early stages of thin film formation 45

3.2.2. 3D thin film growth 48

viii

3.2.3. Microstructure of thin film 50

3.3. Influence of energetic species on thin film properties during sputtering 53

3.4. Formation and evolution of stresses in thin films 56

References 60

4. Zirconia thin film deposition and characterization 63

4.1 Modelling and computational details 64

4.2 Thin film deposition and process monitoring 65

4.2.1 Voltage feedback control unit 68

4.3 Film characterization tools 70

4.3.1 X-ray diffraction (XRD) 70

4.3.1.1 Bragg-Brentano (θ-2θ) mode 71

4.3.1.2 Grazing incidence XRD (GIXRD) mode 72

4.3.2 Transmission electron microscopy (TEM) 72

4.3.3 Secondary electron microscopy (SEM) 73

4.3.4 Chemical composition of zirconia thin films 74

4.3.4.1 Rutherford backscattering spectroscopy (RBS) 74

4.3.4.2 Nuclear reaction analysis 75

4.3.5 Electrochemical Impedance spectroscopy (EIS) 76

4.3.6 Photoluminescence (PL) 77

4.3.7 Heat flux microsensor 79

References 81

5. Influence of oxygen vacancies on the phase constitution of zirconia thin films 83

5.1 Phase stability of oxygen deficient zirconia; quantum chemistry based DFT calculations84

5.2 Synthesis of ZrO2-x by reactive magnetron sputtering 86

5.3 Conclusion 90

References 91

6. Influence of the deposition parameters on the phase constitution of oxygen vacancy doped

thin films 93

6.1 Influence of pressure and discharge current 94

6.1.1 Experimental details 94

6.1.2 Results and discussion 95

6.1.2.1 Evolution of the phase constitution as a function of pressure and discharge current

95

6.1.2.2 Normalized energy flux measurements 97

6.2 Influence of film thickness on the crystal structure of OVSZ films 102

ix

6.2.1 Experimental details 102

6.2.2 Results and discussion 102

6.2.2.1 Evolution of XRD diffractograms as a function of film thickness 102

6.2.2.2 Analysis of film cross-section by SEM and HRTEM 104

6.3 Conclusions 108

References 110

7. Thermal stability of OVSZ thin films 111

7.1 Experimental details 112

7.2 Results and discussion 113

7.3 Conclusion 124

References 126

8. Ionic conductivity of OVSZ thin films 127



8.3 Conclusion 137

References 138

9. Optical properties of OVSZ thin films 139



9.3 Conclusion 146

References 148

Outlook 149

Appendix 154

1

1. Introduction 1.1. Background

Thin films are layers of materials whose thickness ranges from a few monolayers to few

micrometers. These films sometimes exhibit unique properties that cannot be observed in

bulk materials. The history of making thin films dates back to the metal ages and especially

to the ancient craft of gold beating. In ancient times, i.e. more or less 5000 years ago, the

Egyptians were the first to practice this art [1] to make decorative gold leafs. Today thin films

are still being used for many decorative proposes but also to enhance the surface properties

of materials as well as to build functional devices in various application fields such as optics,

mechanics, microelectronics, sensors, energy production such as solar energy, solid oxide

fuel cells.

Obviously, the properties of a film and thereby its area of application is mainly determined

by the elemental composition of the material. But, besides the elemental composition of the

coating, the spatial arrangement of the film forming atoms, i.e. the crystal structure, is also a

decisive factor [2,3]. One example of such is zirconia (ZrO2). Zirconia is a polymorphous

material which exists in three crystallographic phases at atmospheric pressure: (i) the

monoclinic phase (m, space group P21/c) stable up to ~ 1205 °C; (ii) the tetragonal phase (t,

space group P42/nmc) appears from ~ 1205 °C to 2377 °C; and finally (iii) the cubic phase

(c, space group Fm-3m) from 2377 °C to 2710 °C (melting temperature)[4]. Since ZrO2

2

exhibits high chemical stability, high hardness[5], high dielectric constant[6] and prominent

optical properties[7], ZrO2 films have been exploited for a broad range of applications e.g.,

medical applications[8,9], wear resistant coatings[10] and for thermal barrier coatings

(TBCs) [11,12]. However, in the case of pure zirconia, it is not possible to exploit most of

the above mentioned applications as this is restricted by the change in volume of the zirconia-

based components (~5 vol.%) due to the phase transformation upon heating and cooling of

the device, which ultimately leads to the deterioration of the device components [13,14]. To

overcome this situation, the understanding of the synthesis process and the origin of the

properties of interest must be unraveled. An example of the latter is the stabilization of the

high temperature c-phase of zirconia at room temperature with the help of aliovalent dopants

(e.g. Y3+ or Mg2+)[15]. By adding around 12 mol. % of yttria (Y2O3), the c-phase of zirconia

is found to be stabilized at room temperature. This material is known as yttria-stabilized

zirconia (YSZ)[15]. In YSZ, some zirconium (Zr4+) ions are replaced by yttrium (Y3+) so that

to maintain the charge neutrality, for two substituting yttrium cations, one oxygen vacancy is

created. This makes YSZ not only useful for TBCs but also as an electrolyte membrane in

solid oxide fuel cells (SOFC)[16–18] and in oxygen sensors[19] because of its very good

ionic conductivity in between 600 °C – 1000 °C [20].

However, it has also been found that the doping by aliovalent cations leads to the perturbation

of the periodic potential of the oxide-ion array, which results in higher energy barrier for O2-

ions during their diffusion to a vacant site in the solid as compared to intrinsic vacancy-doped

oxides[21]. Therefore, to stabilize high temperature c-phases of zirconia at room temperature

without any doping of yttria, an intense research has been developed during the last 15 years

using various thin film deposition techniques. Among the various thin films growth methods,

3

the most common are electroplating[22], chemical vapor deposition (CVD)[23] and physical

vapor deposition (PVD)[24].

Electroplating is one of the simplest thin film deposition techniques and is very common in

the industry due to its cost effectiveness. In such technique, the object to be coated (the

substrate) is immersed into the solution in which the metal is dissolved. Then negative

potential is applied to the object. In this way, the positive metal ions are attracted to the

surface of the object and a coating is formed. Unfortunately, it’s hard to produce high quality,

dense and defect-free, coatings by such method. This technology also produces a

considerable amount of hazardous wastes.

Chemical vapor deposition (CVD) is a family of processes in which volatile gasses

(precursors) are let into the deposition chamber and the coating is grown on a heated substrate

by a chemical reaction(s) occurring on or in its vicinity. This technique is also widespread in

the industry because of its high deposition rates. It also provides homogeneous films on

complex shaped surfaces and it offers the possibility to deposit a great variety of films, from

metals to organic compounds, with a high control on film composition. The major drawback

of CVD is that most of the time very high temperatures are needed to make the different

constituents react in the gas phase or on the substrate. This means that CVD cannot be applied

to temperature sensitive substrates (e.g. plastic) and also in areas where the thermal expansion

coefficients of films and substrates are different e.g. on steel and electrical components. In

some cases, such limitations can be (partially) overcome by using a plasma to activate the

process. In this case, the process is known as plasma-enhanced CVD.

Physical vapor deposition (PVD) relates to a family of processes in which the film is

deposited on a substrate by the condensation of a vaporized material generated from a solid

4

or liquid source (called target). One such method of PVD is DC sputter deposition. In DC

sputter deposition, the target is bombarded by ions generated in a plasma. The bombardment

of the target induces the ejection (sputtering) of the atoms from the target surface. These

atoms are then transported to the substrate where they condensate and form a film. This

process is generally based on the diode configuration with facing electrodes constituting the

anode and cathode, the plasma being located in between. In such an arrangement, the cathode

plays the role of the target to be sputtered. The disadvantage of DC sputter deposition is that

the plasma is not confined efficiently close to the sputtering region[24] and require sputtering

pressure from several tenths of mTorr to 100 mTorr. One way to confine plasma near the

target and to lower the sputtering pressure is by placing a pair of concentric, permanent,

magnets behind the target. This method is known as DC magnetron sputtering. The

advantages of this configuration are the reduced background gas pressure, the very good

quality of the films, and a better control over the deposition parameters. In such configuration

one can deposit elemental, alloyed, or compound films depending on the target composition.

Compound film materials such as Metal-Nitrides or Metal-Oxides are also synthesized by

adding a reactive gas such as N2 or O2 to the argon carrier gas, in the deposition chamber. In

this case, the technique is called DC reactive magnetron sputtering. The deposition of zirconia

(ZrO2) thin films can be achieved using such kind of process.

In the field of thin film synthesis, in particular in the case of PVD and cold-plasma based

magnetron sputtering methods, several mechanisms have been proposed for the stabilization

of the cubic phase of zirconia without the use of any aliovalent ions. The stabilization is either

related to the grain size[13,25], the energy input during film growth[26], to stresses in the

film[27,28] or to the presence of O vacancies (Vo) or N atoms in the lattice[29]. But it appears

5

that a consensus over what is the mechanism(s) responsible for the stabilization in c-phase

formation has not been reached so far.

1.2. Aim of the research work and strategy

The aim of this thesis work is to understand what drives the synthesis and stabilization of the

high temperature c-phase of zirconia at room temperature. For this purpose, the impact of

oxygen vacancy doping on the phase formation will be first investigated using quantum

chemistry based calculations. Then, a well optimized deposition procedure, based on reactive

magnetron sputtering, will be utilized to synthesize vacancy doped zirconia films. The

deposition parameters (i.e. pressure, discharge current leading to the change in energy

deposited during the film growth) will be varied systematically in order to determine their

influence on the c-phase formation, without doping of aliovalent atoms. In the last step,

thermal stability, the ionic conduction and optical characteristics of the developed material

will be measured and discussed.

This manuscript is organized as follows. Following the introductory chapter, a general

overview of zirconia crystal structure, c-phase stabilization mechanisms, zirconia properties

and applications is given in the chapter 2. In Chapter 3, an overview of the reactive magnetron

sputtering thin film deposition technique is given. In chapter 4, details about the theoretical

calculations and the deposition and film characterization experiments are given. Moreover,

the way that the oxygen vacancy doped zirconia thin films are obtained, are presented in

details. The chapter 5 describes the influence of O vacancies on the phase constitution of

zirconia and the chapter 6 describes influence of deposition parameters on the phase

formation of zirconia thin films and the influence of film thickness. Chapter 7 consists on the

thermal stability of oxygen vacancy stabilized cubic zirconia (OVSZ) thin films. Finally, the

6

properties of the OVSZ films are investigated. Ionic conductivity and photoluminescence

measurements are presented in chapter 8 and 9, respectively.

7

References [1] E.D. Nicholson, The ancient craft of gold beating, Gold Bull. 12 (1979) 161–166. doi:10.1007/BF03215119.

[2] A. Van De Walle, A complete representation of structure-property relationships in crystals., Nat. Mater. 7 (2008)

455–458. doi:10.1038/nmat2200.

[3] G.B. Olson, Computational Design of Hierarchically Structured Materials, Science (80-. ). 277 (1997) 1237–

1242. doi:10.1126/science.277.5330.1237.

[4] J.P. Abriata, J. Garcés, R. Versaci, The O-Zr (Oxygen-Zirconium) system, Bull. Alloy Phase Diagrams. 7 (1986)

116–124. doi:10.1007/BF02881546.

[5] R.C. Garvie, R.H. Hannink, R.T. Pascoe, Ceramic steels?, Nature. 258 (1975) 703.

http://dx.doi.org/10.1038/258703a0.

[6] G.D. Wilk, R.M. Wallace, J.M. Anthony, High-k gate dielectrics: current status and materials properties

considerations, J. Appl. Phys. 89 (2001) 5243. http://dx.doi.org/10.1063/1.1361065.

[7] Q. Zhang, J. Shen, J. Wang, G. Wu, L. Chen, Sol-gel derived ZrO 2 -SiO 2 highly reflective coatings, Int. J.

Inorg. Mater. 2 (2000) 319. http://dx.doi.org/10.1016/S1466-6049(00)00037-4.

[8] F. Namavar, J.D. Jackson, J.G. Sharp, E.E. Mann, K.E. Bayles, Searching for smart durable coatings to promote

bone marrow stromal cell growth while preventing biofilm formation, Biofilm-Material Interact. Tools, Technol.

Oppor. (2007) 904–954.

[9] P.F. Manicone, P. Rossi Iommetti, L. Raffaelli, An overview of zirconia ceramics: basic properties and clinical

applications., J. Dent. 35 (2007) 819–26. doi:10.1016/j.jdent.2007.07.008.

[10] X. Zhou, I. Balachov, D.D. Macdonald, The effect of dielectric coatings on IGSCC in sensitized type 304 SS in

high temperature dilute sodium sulfate solution, Corros. Sci. 40 (1998) 1349. http://dx.doi.org/10.1016/S0010-

938X(98)00018-3.

[11] D.R. Clarke, S.R. Phillpot, Thermal barrier coating materials, Mater. Today. 8 (2005) 22–29. doi:10.1016/S1369-

7021(05)70934-2.

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Solid State Mater. Sci. 4 (1999) 255. http://dx.doi.org/10.1016/S1359-0286(99)00024-8.

[13] S. Shukla, S. Seal, Mechanisms of room temperature metastable tetragonal phase stabilisation in zirconia, Int.

Mater. Rev. 50 (2005) 20.

[14] J. Chevalier, L. Gremillard, A. V. Virkar, D.R. Clarke, The tetragonal-monoclinic transformation in zirconia:

Lessons learned and future trends, J. Am. Ceram. Soc. 92 (2009) 1901–1920. doi:10.1111/j.1551-

2916.2009.03278.x.

[15] H.G. Scott, Phase relationships in the zirconia-yttria system, J. Mater. Sci. 10 (1975) 1527–1535.

[16] S.M. Haile, Fuel cell materials and components, Acta Mater. 51 (2003) 5981–6000.

doi:10.1016/j.actamat.2003.08.004.

[17] B.C. Steele, A. Heinzel, Materials for fuel-cell technologies., Nature. 414 (2001) 345–352.

doi:10.1038/35104620.

[18] A. Atkinson, S. Barnett, R.J. Gorte, J.T.S. Irvine, A.J. Mcevoy, M. Mogensen, S.C. Singhal, J. Vohs, Advanced

anodes for high-temperature fuel cells, Nat. Mater. 3 (2004) 17–27.

[19] A. Cirera, C. Lpez-Gándara, F.M. Ramos, YSZ-based oxygen sensors and the use of nanomaterials: A review

from classical models to current trends, J. Sensors. 2009 (2009). doi:10.1155/2009/258489.

[20] J. Nowotny, M. Rekas, T. Bak, Defect Chemistry and Defect-Dependent Properties of Undoped and Stabilised

Zirconia. Bulk vs Interface, Key Eng. Mater. 153–154 (1998) 211–240.

doi:10.4028/www.scientific.net/KEM.153-154.211.

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doi:10.1038/35009177.

[22] J.E. Gray, B. Luan, Protective coatings on magnesium and its alloys — a critical review, J. Alloys Compd. 336

(2002) 88–113. doi:10.1016/S0925-8388(01)01899-0.

[23] K.F. Jensen, Chemical Vapor Deposition, Third Edit, Elsevier Ltd., 2010. doi:10.1016/B978-0-8155-2031-

3.00007-7.

[24] D.M. Mattox, Handbook of Physical Vapor Deposition (PVD) Processing, 2nd ed., Elsevier Inc., 2010.

[25] S. Shukla, S. Seal, R. Vanfleet, Sol-Gel Synthesis and Phase Evolution Behavior of Sterically Stabilized, J. Sol-

Gel Sci. Technol. 27 (2003) 119–136.

[26] S. Mraz, J.M. Schneider, Influence of the negative oxygen ions on the structure evolution of transition metal

oxide thin films, J. Appl. Phys. 100 (2006) 23503. doi:10.1063/1.2216354.

[27] J.M. Ngaruiya, O. Kappertz, S.H. Mohamed, M. Wuttig, Structure formation upon reactive direct current

magnetron sputtering of transition metal oxide films, Appl. Phys. Lett. 85 (2004) 748. doi:10.1063/1.1777412.

[28] K. Goedicke, J. Liebig, O. Zywitzki, H. Sahm, Influence of process parameters on the structure and the properties

of ZrO 2 coatings deposited by reactive pulsed magnetron sputtering (PMS), Thin Solid Films. 377–378 (2000)

37–42.

[29] K. Sarakinos, D. Music, S. Mraz, M. to Baben, K. Jiang, F. Nahif, a. Braun, C. Zilkens, S. Konstantinidis, F.

Renaux, D. Cossement, F. Munnik, J.M. Schneider, On the phase formation of sputtered hafnium oxide and

oxynitride films, J. Appl. Phys. 108 (2010) 14904. doi:10.1063/1.3437646.

8

9

2. Zirconia 2.1. The Crystal structure of Zirconia (ZrO2)

Zirconia (ZrO2) is a polymorphous material. At atmospheric pressure it can be found in three

fluorite related crystallographic phases; (i) the monoclinic (m), occurs naturally as a mineral

Baddeleyite and is stable up to ~ 1205 °C; (ii) the tetragonal phase (t) appears from ~ 1205

°C to 2377 °C; and finally (iii) the cubic phase (c) from 2377 °C to 2710 °C (melting

temperature)[1]. The unit cells related to m-, t-, and c-zirconia are presented in Fig.2.1 and

their characteristic crystallographic properties on Table 2.1.

Fig. 2. 1: Crystal structure of Zirconia (a) monoclinic, (b) tetragonal, (c) cubic

Table 2. 1 Crystal structure of pure zirconia (ICDD PDF # 013 0307; 010 070 6627; 00 049 1642).

Temperature

(°C) Phase

Space

group

Space

group number

Cation

coordination number

Cell parameters

a(Å) b(Å) c(Å) α(°) β(°) γ(°)

0-1205 Monoclinic P21/c 14 7 5.1477 5.2030 5.3156 90.00 99.38 90.00

1205-2377 Tetragonal P42/nmc 137 8 3.5948 3.5948 5.1824 90 90 90

2377-2710 Cubic Fm-3m 225 8 5.1280 5.1280 5.1280 90 90 90

10

Zirconia high temperature c-phase in general has superior properties over the room

temperature stable m- and t-phase because of highly symmetric atomic arrangement.

Therefore, for most of the engineering applications, zirconia c-phase is preferred. However,

this is limited by the instability of cubic zirconia at room temperature. Further limitation in

the use of zirconia for high temperature applications comes from the change in its unit cell

volume as a function of temperature [2], as shown in Fig 2.2. It could be observed the unit

cell volume of zirconia first increases up-to 1400 K and then, due to the phase transformation,

the unit cell volume decreases abruptly. This change in volume during the increase in

temperature as well during the cooling results in the formation of cracks in the coating, thus

making it un-useful in such circumstances.

Fig. 2. 2: Variation of the Zirconia unit cell volume of m, t, and c-phase as a function of temperature,

adopted from [2]. The t-phase here is described in terms of face-centered tetragonal cell of unit cell

parameter á. The usual crystallographic has a = a/√2 [2].

2.2. Stabilization strategies for zirconia c-phase

Two strategies are found in the literature regarding the stabilization of the high temperature

cubic phase of zirconia at room temperature. One is to introduce a foreign element (doping)

11

in the zirconia lattice. The second is related to a mechanism that does not involve any doping

process Here (in section 2.2.1 and 2.2.2) both approaches are presented.

2.2.1. Stabilization of Zirconia c-phase by doping

Fluorite oxides (MO2) are the oxides which exhibit cubic crystal structure consisting of a

cubic oxygen lattice with alternate body centers occupied by eight coordinated cations. The

cations are arranged into a face centered cubic structure with the anions occupying the

tetrahedral sites. Usually the size of these tetravalent cation (M) is big enough to sustain the

cubic (fluorite) structure. However, in case of zirconia (ZrO2), the size of Zr4+ is too small to

sustain the cubic structure at low temperatures. Therefore, to stabilize the cubic structure, it

has to be partly substituted with a larger cation than Zr4+. For this purpose, since more than

ninety years, the stabilization of the high temperature c-phase of zirconia at room temperature

is achieved by doping zirconia (ZrO2) by larger cations of lower valence than Zr4+, e.g., by

incorporating Y3+, Ca2+, or Mg2+ in the ZrO2 lattice. This strategy was first reported by Ruff

et al.[3] in 1929. In the beginning, due to the lack of advanced measuring equipment, the

dopant values reported for the stabilization of c-phase were a bit higher than the recently

published values. It has been reported by Gaudon, et al.[4] that by doping around and above

7 mol% of yttria (Y2O3), the c-phase of zirconia can be stabilized at room temperature, as

shown in Fig.2.3. This material is known as yttria stabilized zirconia (YSZ). On the other

hand, doping by 2-7 mol% of yttria, leads to the stabilization of the tetragonal phase also

known as partially stabilized zirconia. Below 2 mol% of yttria doping, the monoclinic phase

is present.

12

Fig. 2. 3: Phase diagram of Zirconia-yttria system. m, t and c represent the monoclinic, tetragonal and

cubic phase respectivley[4].

In YSZ, on addition of Y3+ in the zirconia lattice, Zr4+ cations are replaced by Y3+. To

maintain the charge neutrality, for each two substituting yttrium cations, one oxygen vacancy

is created. This process is summarized by the eq. (1 and 2) using Kröger-Vink notation and

is schematically illustrated in Fig. 2.4.

𝑌2𝑂3 𝑍𝑟𝑂2→ 2𝑌′𝑍𝑟 + 𝑉𝑂

.. + 3𝑂𝑂𝑥 (1)

𝑂𝑂𝑥 ↔ 𝑉𝑂

.. + 2𝑒 +1

2𝑂2 (2)

𝑌′𝑍𝑟 represent the Y in the Zr site with the apparent negative charge, 𝑉𝑂.. is the vacancy in the

oxygen site with double positive charge, 𝑂𝑂𝑥 is the lattice oxygen, i.e., oxygen in the oxygen

site with net charge of zero.

13

Fig. 2. 4: Schematic representation of the insertion of Yttria (Y2O3) in the zirconia (ZrO2) lattice and

the creation of oxygen vacancies. Image adopted from [5].

2.2.2. Stabilization of high temperature phase of Zirconia without doping

In order to stabilize t- and c-phase of zirconia at room temperature, especially in case of

zirconia thin films, several mechanisms have been proposed so far. The stabilization has been

attributed to the grain size [6–8], energy input during the film growth [9–11], O vacancies

and/or incorporation of N atoms in the lattice [12–14] and has been also related to the stresses

in the film [8,9,11]. These mechanisms are briefly presented here.

Grain-size effect

The effect of grain size on stabilizing the zirconia t-phase at room temperature without any

dopants was first reported by Gravie [15] in 1965. In their study they reported that 100% t-

ZrO2 can be achieved at room temperature by having a grain size in the range of 11-17 nm.

However, if the grain size ranges from 17 nm to 30 nm, a mixture of tetragonal and

monoclinic phases was found. It is concluded in their study that a critical size exists for the

14

stabilization of the metastable tetragonal phase, which is found to be 30 nm. Further it has

been also reported by Chen et al. [16] that the tetragonal phase of zirconia can also be

stabilized for temperature in the order of 700 °C (as shown in Fig. 2.5). This is a much lower

temperature range than what is reported on the phase diagram (see Fig. 2. 3) where the

temperature ranges from 1205 °C to 2377 °C. According to these authors, such a lowering of

the temperature at which the t-ZrO2, but also the metastable cubic phase, is obtained is

achieved by applying an external hydrostatic pressure on powdered zirconia. This is due to

the densities of tetragonal, orthorhombic, and cubic crystal structures which are higher than

that of the monoclinic structure. The application of an external hydrostatic pressure converts

the monoclinic structure, which is stable at lower temperatures, into a denser structure.

Further a relation between the grain size and phase transformation temperature was also

found by Garvie et al. [17,18]. Garvie et al. shows the phase transformations temperature

Fig. 2. 5: Schematic of pressure-temperature phase diagram of bulk ZrO2. Image adopted from [19].

15

can be lowered by reducing the grain size[17,18], therefore it’s possible to obtain the

tetragonal and cubic phases at room temperature instead of Orthorhombic-I, -II by applying

an external pressure. Shukla et al.[7] have reported in their study that an external hydrostatic

pressure is required to stabilize the tetragonal/cubic phase at room temperature for zirconia

grain size greater than 10 nm. However, in the case of a grain size less than 10 nm, there is

no need to apply external hydrostatic pressure. Therefore, there must be sufficient hydrostatic

pressure acting inside the grain to stabilize the tetragonal and cubic phases at room

temperature in this situation. In the case of liquid particle of radius (r); the magnitude of the

hydrostatic pressure (∆P) can be calculated from Gibbs-Thompson equation

∆𝑃 =2𝛾

𝑟(3)

where γ is the surface tension. In the case of solids, the surface tension has to be replaced by

the surface stress f, which is given by the expression (4) where ε is the strain.

𝑓 = 𝛾 +𝑑𝛾

𝑑𝜀(4)

Since in solids, internal hydrostatic pressure is due to the surface stress which is of the order

of surface energy. Therefore, for spherical isotropic particles dγ/dε can be neglected. Thus

the surface tension γ can be substituted by surface stress f in in equation (3)

∆𝑃 =2𝑓

𝑟(5)

The magnitude of f has been estimated by Winterer et al. for nano-crystallite zirconia powder

with a grain size ranging from 5-30 nm and is equal to 5 Nm-1[20]. Substituting this value in

equation 5 and calculating the internal hydrostatic pressure for grains of 10 nm and below

reveals that the internal hydrostatic pressure increases with the decrease in grain size and is

as high as 25 GPa for a grain-size of 1 nm, (Fig. 2.6). Further, Winterer et al.[20] estimated

16

the grain size for the stabilization of tetragonal phase. This critical grain size is equal to 8-12

nm. The later value indicates that the internal hydrostatic pressure should be about 2.5 GPa.

On the other hand, to stabilize cubic phase of zirconia, a slightly larger internal hydrostatic

pressure (~30 GPa) is required. Nitsche et al.[21,22] also synthesized nanosized zirconia

powder characterized by grain size from 5-30 nm using a gas condensation technique.

However, from HRTEM analysis, they found that the size of the zirconia grains lies in the

range of 7-32 nm and the grains exhibit a core-shell morphology with the tetragonal phase as

the core and the monoclinic phase as the shell.

Fig. 2. 6: Variation in internal hydrostatic pressure as a function of Nano-crystallite (grain) size,

calculated using equation (5), image adopted from [7].

Influence of energy deposition and stresses on phase stabilization

In a plasma based process, a film grown on the substrate surface is subjected to intense

particle bombardment. The impinging species can be among others, plasma ions (e.g. Ar+,

O2-, Zr+), electrons, photons, condensing metal atoms, neutral particles [23]. The flux of each

of these particles will influence to amount of energy delivered to the film during growth. The

17

influence of energy input on stabilizing high temperature cubic phase of zirconia was first

noticed by Goedicke et al.[9] in 2000, for films of 100-150 nm thickness synthesized by

reactive pulsed magnetron sputtering (PMS) at various deposition pressure. Goedicke et al.

varied the working pressure from 0.3 Pa to 3.5 Pa as well as the target to substrate distance.

During the deposition, they did not supply any intentional heating to the substrate. However,

they observed an increase of maximum 30 °C in the substrate temperature during the

deposition by using thermos strips. In their study, when the films were deposited at low

pressure i.e. 0.3 Pa, films exhibited pure monoclinic phase. On the other hand, when films

were deposited at 3.5 Pa, they exhibited pure cubic phase as shown in Fig. 2.7. The reason

Goedicke et al. propose for such a behavior is related to the energy of the condensing

particles. They suggested that higher sputtering pressure reduces the mean free path of the

sputtered particles and therefore the mobility of the depositing species on the substrate

surface. They further relate lower mobility to microporosity of the film and therefore reduced

hardness and lower compressive stresses in the film. Goedicke et al. also measured residual

stresses in their study and found that, indeed, films deposited at low sputtering pressure (0.3

Pa) exhibit high compressive stresses (~ 1800 MPa) while the films deposited at higher

sputtering pressure (3.5 Pa) exhibit low tensile stresses (~139 MPa).

18

Fig. 2. 7: Diffractograms of ZrO2 thin films deposited by PMS at sputtering pressure of (a) 0.3 Pa,

(b) 3.5 Pa. Image adopted from [9].

In reactive magnetron sputtering having O2 as a reactive gas, an important feature is the

emission of negatively charged oxygen ions O- from the oxidized part of the target

surface[11,24–26]. These O- ions are accelerated in the cathode sheath. They bombard the

growing film with energies higher than the other depositing plasma species, in the range of

several hundreds of eV [26]. The energy of these O- ions is correlated to the magnitude of the

target voltage, while the number of these O- ions is determined by portion of the target

covered with the oxide layer. In several studies, the influence of these O- ions on the

19

formation of crystalline structures has been reported. For example, in 2006, Mraz et al.[27]

studied the influence of the energy brought by these oxygen ions emitted from the oxidized

target during the growth of transition metal oxides (Nb, Ta, Zr, and Hf) films by reactive

magnetron. These authors proposed that the evolution of the crystalline structure of transition

metal oxide thin films may depend on the presence of O− ion bombardment induced adatom

mobility. Further, Ngaruiya et al. [11] also studied the structure formation of various

transition metal oxides of group 4 (Ti, Zr, Hf), 5 (V, Nb, Ta) and 6 (Mo, W) deposited by

reactive magnetron sputtering at 6 mTorr. Ngaruiya et al. found that Zr and Hf-based

sputtering processes allow for the formation of the monoclinic phases of their respective

oxides. While the other transition metals from group 5 (V, Nb, Ta) and 6 (Mo, W) form

amorphous films. For the zirconium target, Ngaruiya et al. observed that, with the increase

in oxygen flow, the films show a decrease in compressive stresses up to a critical limit which

marks the onset of complete target oxidation. At complete target oxidation, film stresses

increased abruptly (up to (-1500 MPa) and these fully oxidized films were found to be in the

monoclinic phase. They attributed the generated compressive stresses in the film to the

energetic bombardment of oxygen negative ions as result of target oxidation. The latter is

monitored by observing the target voltage as a function of the oxygen flow introduced in the

chamber (Fig. 2. 8). For group 4 oxides (Ti, Zr, Hf), it is observed that the cathode potentials

are moderate as compared to group 5 and 6, which result in providing moderate O- ion energy

for film crystallization and atomic arrangement that causes stresses build up. Finally, it is

concluded that the high flux of low energy oxygen negative ions emitted from the Zr target

accounts for the crystallization and the stress build up in the zirconium oxide films. On the

other hand, for group 5 (V, Nb, Ta) and 6 (Mo, W) oxides, it is proposed that the oxygen ions

get accelerated by the high cathode potential and bombard the growing film. To the contrary

20

of the group 4 transition metals, in this case, the swift bombardment results in the relaxation

of stresses via thermal spikes and in the production of an amorphous structure.

Fig. 2. 8: Variation in Zr target voltage and film stresses as a function of oxygen flow [11].

In 2006, Severin et al.[28] controlled the target oxidation state and therefore the emission of

the fast O- ions by adding N2 to the sputtering process. Later in 2008, Severin et al.[10]

reported the stabilization of the high temperature cubic phase of zirconia films at room

temperature by using such N2 addition strategy (Fig. 2.9). When nitrogen is not added in the

sputtering process and zirconium is sputtered in an argon/oxygen atmosphere, the deposited

films are mainly amorphous and monoclinic. There is only a very small window around 1.7

sccm of oxygen flow where the cubic zirconia films are obtained. On addition of 0.75 sccm

nitrogen flow, the cubic phase zirconia thin films are deposited at lower partial pressure of

21

O2 and the operating window is broadened. This window is further enlarged by increasing

the nitrogen flow to 1.5 sccm and lowering the oxygen flow.

Fig. 2. 9: Phase diagram of reactively sputtered ZrOxNy films for various oxygen and nitrogen flows

as reported by Severin et al.[10].

In 2010 Sarakinos et al.[13] demonstrated for HfO2 (which is isostructural to ZrO2) that it’s

not the bombardment of O- ions which governs the cubic or tetragonal phase formation but

the incorporation of O vacancies and/or the substitution of O atoms by N atoms in the

nonmetal sub-lattice. Sarakinos et al. used the same dc-reactive magnetron sputtering and

sputtering ambient as the one reported by Severin et al. [28]. To prevent the fast O- ions

emitted from the target (i.e. the racetrack) to reach the film, they mounted a Cu ring above

the sputtering target, as presented in Fig. 2.10. It should be noted here that Sarakinos et al.

used a high power pulse magnetron sputtering discharge (HPPMS) to grow the films. This

technique is known to enhance ion bombardment during film growth.

22

Fig. 2. 10: Schematic of the strategy employed by Sarakinos et al to block the bombardment of O-

ions on the growing film. Image adopted from [13].

Influence of oxygen vacancies on zirconia c-phase stabilization

The role of oxygen vacancies in the stabilization of zirconia tetragonal and cubic phase has

been also studied theoretically by Fabris et al. [12]. for a cell of 96 atoms (32 Zr, 64 O) these

authors demonstrated that by incorporating 1 (equivalent to incorporating ~3.2 % mol Y2O3)

and 4 (equivalent to incorporating ~14.4 % mol Y2O3) oxygen vacancies in the zirconia lattice

i.e. by removing oxygen atom(s) from the cell, that the presence of oxygen vacancies is

responsible for the stabilization of the high temperature tetragonal and cubic phase of

zirconia. However, they also suggested that a such stabilization procedure i.e. by doping the

zirconia crystal with oxygen vacancies, may be achieved theoretically only.

2.3 Properties and applications of zirconia

Zirconia is an extremely versatile material finding its applications in medical field[29–32],

as decorative coatings [33–35], cutting blades [7], thermal barrier coatings [36–38], in

23

oxygen sensors [39–41], solid oxide fuel cells [41–45] and in other high temperature

applications because of its superior chemical stability, optical, mechanical, thermal and ionic

properties. Therefore, a brief overview of these mechanical, thermal, ionic, and optical

properties is given in the next sections.

2.3.1 Mechanical properties of zirconia

The elastic properties of pure bulk zirconia have been studied by Chan et al. [46] as a function

of the crystal structure of the material. It is reported that the bulk moduli of monoclinic and

tetragonal zirconia hovers around 150-200 GPa while cubic zirconia has a bulk modulus

around 171-288 GPa. On the other hand, high pressure phases i.e. orthorhombic-I and

orthorhombic-II phases, have values around 224-273 GPa and 254-444 GPa, respectively.

Further, the monoclinic phase of zirconia has hardness about 9.2 GPa[47] for samples with a

density >98% and 4.1-5.2 GPa[48] for samples with a density >95%, whereas hardness of

amorphous zirconia vary between 5 and 25 GPa[49]. The hardness values have been observed

to increase with the addition of yttria in the zirconia lattice, e.g. the hardness approached 11

GPa with a doping of 1.5 mol% of yttria [48]. The addition of larger amounts of yttria i.e. 8

mol.% of yttria leads to a further increase in the hardness which reaches a value of 15 GPa

[50].

As can be seen from the above mentioned values of elastic modulus and hardness, tetragonal

and cubic zirconia have superior mechanical properties than monoclinic. This is the reason

why most of the mechanical engineering applications make use of tetragonal and/or cubic

phases by stabilizing them at room temperature with the help of dopants.

24

2.3.2 Thermal properties of zirconia

Zirconia in its monoclinic phase exhibits a thermal conductivity of 7.2 W/m.K [51] at room

temperature. This value decreases to 2.5 W/m.K at 1100 °C (see Fig. 2.11). This makes this

material very interesting for thermal barrier coatings (TBCs) [36]. However, an increase in

temperature leads to the phase transformation (shown in Fig. 2.2) and thus causes

delamination and cracks in the coating. This is avoided by doping the zirconia lattice with

yttria [36], i.e. by stabilizing the high temperature cubic phase of zirconia. Such stabilization

not only help to avoid phase transformation but also lowers the thermal conductivity of YSZ,

e.g., from 1.42 W/m.K (at room temperature) to 1.35 W/m.K at 1200 °C as c-phase of zirconia

has lower thermal conductivity as compared to m-zirconia.

Fig. 2 11: Thermal conductivity of pure zirconia (ZrO2) [51].

2.3.3 Ionic conduction of stabilized zirconia

Stabilization of zirconia at room temperature with the help of dopants not only make it useful

for mechanical and thermal barrier applications but also as an ionic conductor (electrolyte)

to be incorporated in solid oxide fuel cells (SOFCs) and oxygen sensors. It is because, on

25

addition of aliovalent ions in the zirconia lattice, oxygen vacancies are created to maintain

the charge neutrality.

It was first expected that the ionic conductivity will increase with the increase in O vacancy

concentration i.e. by increasing dopant content (Y2O3). However, later it was observed the

maximum ionic conductivity in yttria-stabilized zirconia (YSZ) occur at 7 – 9 mol% Y2O3 at

327 – 1227 °C [52,53]. On the other hand, higher amount of Y2O3 was found to lower the

mobility of O vacancy by increasing the diffusion energy across an Y–Y common edge as

compared to the diffusion across one with a Zr–Y common edge [52]. The formation of

oxygen vacancies allow oxygen ion O2- migration[43,54], as schematically represented in

Fig. 2.12. YSZ is one of the most frequently used electrolyte materials in SOFCs and oxygen

sensors because of its great availability, cost effectiveness, stability, and high ionic

conductivity.

Fig. 2. 12: Schematic of oxygen vacancy migration in YSZ, Image adopted from [5].

Stabilized zirconia and fuel cells for energy production

A Fuel cell produces electricity from the electrochemical combination of a fuel with an

oxidant. A fuel cell consists of an anode and a cathode separated by an electrolyte. The fuel

(e.g., hydrogen) is fed to the anode where it is oxidized and electrons are released to the

external (outer) circuit. While the oxidant (e.g., oxygen) is fed to the cathode where it is

26

reduced and electrons are accepted from the external circuit. So, electrons flow from the

anode to the cathode through an external circuit and produce an electric current. The

electrolyte conducts ions between the anode and cathode. The key feature of fuel cells is their

high energy conversion efficiency because the fuel cell directly converts chemical energy

into electrical energy. Moreover, fuel cells also offer many advantages over traditional energy

conversion methods e.g., significant higher conversion efficiency, modular construction, high

efficiency at part load, minimal siting restriction, potential for cogeneration, and the most

important, much lower production of pollutants [43].

The principle of fuel cell was first reported in 1839 by Sir William Grove [55]. His fuel cell

used diluted sulfuric acid as an electrolyte and operated at room temperature. The discovery

of solid-oxide electrolytes came much later, in 1899, by Nernst [56] and the operation of the

first fuel cell containing solid-oxide electrolyte (also called solid oxide fuel cell) at 1000 °C

was reported by Baur and Preis in 1937 [57]. Since then, SOFC technology has developed a

lot. Mutlikilowatt fuel cells based on stabilized zirconia electrolyte have been operated for

thousands of hours and have shown good performance[43]. However, the conductivity

requirement for the electrolyte determines the operating temperature of the SOFC e.g., the

ionic conductivity of 8 mol% yttria stabilized zirconia equals 1x10−5 S/cm at 500 °C but

ramps up to 4x10−3 S/cm at 900 °C [43,44,54,58]. This situation requires to the use of high

operating temperatures for SOFCs, typically about 1000 °C [43]. For the device to operate at

such high temperatures, the device components must need to meet certain requirements. For

example, each component of device must have the proper stability (chemical, phase,

morphological, and dimensional) in oxidizing and/or reducing environments, chemical

compatibility with other device components, and proper conductivity. In addition to these,

the components of device must have similar coefficients of thermal expansion to avoid

27

separation or cracking during fabrication and operation. Beside the above mentioned

requirements, the electrolyte and interconnect must be dense to prevent the gas mixing. The

electrolyte used in the SOFCs are normally of thickness ranging several micrometers (thick

films). However, the use of thicker films leads to the high ohmic losses. Therefore, to

overcome such high ohmic losses high operating temperatures are used. An alternate way to

overcome this limitation is the use of thin films (tens of nm thick) e.g., superionic

conductivity i.e. 1S/cm has been reported for 58 nm thick YSZ films in the temperature range

150-500 °C [59] This will not only help to lower the operating temperature and ohmic losses

but will also help to lower the manufacturing cost of the SOFCs by allowing the use of low

cost device components.

2.3.4 Optical properties of Zirconia

Zirconia has also very good optical properties. The refractive index typically equals 2.1-2.2

[9,60,61]. Zirconia also presents a large optical band gap of 5.1-6 eV[62–66] and is therefore

transparent in the visible range. Zirconia is also a high K dielectric (K=25) [67] material,

making zirconia a good candidate in the silicon microelectronics element base applications.

Further, zirconia is a prospective candidate for the role of an active medium in the next

generation Resistive Random Access Memory (ReRAM) [68]. The optical and transport

properties of zirconia, as for other high-K dielectrics, are determined by the presence of

defects in the structure. According to the literature, a blue luminescence band with an energy

of about 2.5-2.8eV is observed in zirconia [67,69–73]. However, the blue luminescence is

attributed to the presence of defects and impurities in the zirconia lattice but the origin of

these luminescence is unclear.

28

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31

3. Thin film growth

Thin films can be deposited in a number of ways. In this regard, methods commonly used to

synthesize tetragonal and cubic zirconia thin films are presented in Fig. 3. 1. As discussed in

section 1.1, zirconia thin film deposition can be divided in two major disciplines, i) chemical

deposition methods and ii) physical vapor deposition methods (PVD). Chemical deposition

method further can be divided into chemical solution based synthesis and chemical vapor

deposition (CVD) method. In chemical solution based method, precursor solution is

employed on the substrate and later processed to have a thin film, it includes Sol-gel method

[1–4] and Spray pyrolysis [5–7] methods. On the other hand, in chemical vapor deposition

method, volatile gasses (precursors) are let into the deposition chamber, which is under

vacuum and the zirconia coating is grown on a heated substrate by a chemical reaction

occurring on or in its vicinity. CVD itself is the parent of a family of film growth processes.

The most common sub-CVD methods for zirconia thin film deposition are atomic layer

deposition [8–12] and metal organic chemical vapor deposition (MOCVD) [13–17]. Like

CVD, the term PVD also includes a family of techniques. During physical vapor deposition,

a vapor of the film forming species is created from a solid, or sometimes, from a liquid source,

by physical means. The source of material is usually named the target and the vapor of

depositing species can be obtained e.g. through heating or as a result of the interaction of fast

ions or a LASER beam with the target surface. The vapor is then transported through the gas

32

phase to the substrate, where it condensate and form the film. The common methods among

PVD to deposit zirconia thin films include cathodic arc evaporation [18], pulsed laser

deposition [19–22], ion beam assisted deposition [23–25] and sputtering [26–36].

Fig. 3. 1: Flowchart summarizing the most common deposition methods used to deposit zirconia

thin films.

3.1. Sputtering

Sputtering is a widely used method not only to deposit binary compound materials such as

Metal-Oxides or Metal-Nitrides thin films but also to deposit elemental, alloy, mixture or a

compound film depending on the target composition. Further the choice of sputtering comes

from its conceptual simplicity as well as from its easy scalability. Sputtering is by definition

the ejection (removal) of atoms from the target material which is bombarded by energetic

species. The bombarding species are ions or fast neutrals extracted from the plasma but

(positive) ions are usually considered as the most important particle. Depending on the energy

of the incident/bombarding ion, the sputtering process can be divided into three regimes; i)

single knock-on (low energy, 10-30 eV), ii) linear cascade (moderate energy, 100 eV-10 keV)

33

and iii) spike (high energy, >10 keV ), as shown in Fig. 3.2. In a single knock-on event, in

the low energy regime, a small fraction of the target atoms is set into motion by interaction

of incident ion with the surface atoms. The bombarding ion provides enough energy to the

surface atoms to overcome the surface binding energy and to sputter out from the target

surface. While in linear cascade regime, the incident particle goes under a series of collisions

with several target atoms. This situation triggers the displacement of several target atoms

from their sites, which go in collision with other surface atoms (linear cascade) and thus

sputtering them out. In the spike regime, the incident ion carries a substantially high amount

of energy which provide enough energy to all the atoms along its path to overcome their

binding energy, causing a higher spatial density of moving atoms as compared to linear

cascade regime. In the spike regime the density of recoil atoms is so high that the majority of

atoms with in the spike volume are in motion and the region of collisions become so dense

that multiple collisions occur simultaneously. In this case the process becomes a complicated

process of many-body interactions between hundreds and tens of thousands of atoms, which

cannot be treated with the binary collisional process (as in the case of linear cascade).

The number of atoms ejected out from the target per incident ion is referred as the sputtering

yield. The latter depends on the incident ion energy, the angle of incidence, the surface

binding energy of the target material, and on the mass of the incident ion and the target atom

[37]. Beside the ejection of atoms from the target surface during sputtering, some additional

phenomena are also observed. Such phenomena include backscattering of the incident

species, change in the surface structure and morphology, as well as the emission of electrons

from the target surface and the production of phonons[38]. The electrons emitted from the

target by the ion bombardment are known as secondary electrons. The emission of secondary

electrons is of particularly important in the case of plasma-based sputtering methods. Indeed,

34

the secondary electron emission results in the generation of new ions through inelastic

collisions with atoms in the surrounding gas phase, these newly formed ions will in turn

bombard the target. This situation enables the continuous emission of secondary electrons

and gas atom ionization which allow the formation of a quasi self-sustained glow discharge.

Glow discharges are discussed in the next section.

Fig. 3. 2: Sputtering regimes (a) Single knock-on (low energy) (b) Linear cascade (c) Spike (high

energy) [37].

3.1.1. Glow discharge

The sputtering process is generally based on a diode configuration with facing electrodes,

where the cathode is the target to be sputtered. To realize sputtering, high-purity noble gas

(typically Ar) is introduced in the continuously pumped chamber. The working pressure

ranges from a few tenths of mTorr5 up to 100’s of mTorr. Then a constant negative potential

is applied to the target using an electric power supply. This result in the acceleration of the

few electrons already present in the low pressure gas. These initial electrons-ions pairs are

generated by external sources such as X-Rays or UV photons coming from the surroundings

(e.g. cosmic rays). The primary electrons, accelerated by the electric field of the cathode

(typically hundreds of volts), promote a wave of ionization in the noble. The positive ions

5 1 mTorr = 0.133 Pa

35

(Ar+) produced during this ionization cascade can be in turn accelerated towards the

negatively biased sputter target. The ion bombardment results in the removal of target atoms

(sputtering) as well as in the emission of secondary electrons as described in the previous

section. These fast secondary electrons allow to sustain the ionization of the gas phase. As

the steady state is reached, a partially ionized gas consisting of energetic particles i.e. ions,

electrons, photons and a much larger number of gas neutrals is obtained. This particular

medium is called glow discharge or plasma[39]. The term plasma was first used by Irving

Langmuir in order to describe an ionized gas in 1927[40].

The generation and stabilization of the plasma is the core of the sputtering processes. For this

purpose, collisions in the gas phase are essential. Since the gas phase is a combination of

ions, electrons, neutrals and molecules, therefore ideally one should consider the interaction

between all possible pair of permutations. These collisions can be divided into elastic and

inelastic collisions. In elastic collisions, the kinetic energy of the particle is transferred to the

target particle and the target particle is accelerated i.e. the internal energy of the target particle

remains unaffected. In inelastic collisions, part of the kinetic energy of the missile is

transferred to the target particle but in this case, the internal energy of the target particle

increases. Hence, the target particle can be excited, ionized, or dissociated (if the target

particle is a molecule).

In low pressure plasmas such as those utilized for sputtering, the collisions involving

electrons usually dominate the ionization processes[39]. Ionization by electron impact is

allowed if the electron has at least a high enough kinetic energy i.e. above the ionization

potential of the target atom or molecule. Actually, every collisions process occurring in the

gas phase is characterized by a certain probability which is determined by the so-called

36

collision cross-section. Cross sections are energy dependent quantities[39]. Other ionization

processes observed in the plasma phase include ion-neutral and metastable-neutral collisions.

The counter part of the ionization is electron-ion recombination. Due to the recombination

processes which mainly happens at the chamber walls, an external source is always needed

for the generation of new electrons to sustain the plasma. This is the role of the negative

potential applied to the target as described in the beginning of this section. Further, a less

dramatic transfer of the energy would cause excitation, where the bound electrons of the

target atom/molecule jump to a higher energy level. As the lifetime of the excited state is

finite, the excited electron decays to a lower energy state after that finite time by releasing

the excess energy in the form of a photon. The emitted photon may have energies

corresponding to the visible spectral range. Consequently, the plasma glows.

Formation of sheath

Plasma in general is a quasi-neutral i.e. there is an equal amount of negative and positive

particles (ne≈ni)[39]. This changes dramatically when the plasma is locally disturbed, e.g.

when an electrically isolated (floating) surface is inserted into the plasma. In this situation,

the substrate gets exposed to a flux of electrons, ions and neutral atoms. The flux of impinging

species is, in a first approximation, equal to 𝑛𝑐

4 [39] where n is the density of the

corresponding particles and 𝑐 their average velocity. In the very beginning, since the

electrons are much faster than the other plasma species, because of their much lower mass,

the floating substrate experiences a much higher electron current density as compared to the

ion current density. This result in the buildup of negative net charge on the floating substrate

and hence of a negative potential with respect to the plasma. Consequently, newly arriving

electrons get repelled by this negative potential while positive ions are attracted. This

37

negative potential appearing on the electrically insulated surface is known as the floating

potential (Vf) and the potential of the undisturbed bulk plasma is the plasma potential (Vp).

The plasma potential has the most positive value in the glow discharge. Once electron and

ion fluxes equilibrate, a potential difference equal to |Vp|-|Vf| appears between the plasma

bulk and the substrate surface. As a consequence, only electrons with energy higher than the

potential difference |Vp|-|Vf| can overcome the barrier and a net positive charge appears

around the surface. Locally the neutrality relation (ne≈ni) doesn’t hold anymore and this

space charge region is named sheath [39].

Now let’s assume that the kinetic energy of the electrons obeys the Maxwell-Boltzmann

distribution, then the plasma density in the sheath ��𝑒 can be written as[39]

��𝑒𝑛𝑒= 𝑒𝑥𝑝 − (

𝑒(𝑉𝑝 − 𝑉𝑓)

𝑘𝐵𝑇𝑒) (3.1)

Here ne is the plasma density of undisturbed plasma and Te is the electron temperature. Since

the plasma potential difference |Vp|-|Vf| result in the spatial variation of the potential

(∆𝑉(𝑥))within the sheath and is given by the equation[39],

∆𝑉(𝑥) = (𝑉𝑝 − 𝑉𝑓) exp (−𝑥

𝜆𝐷) (3.2)

where 𝜆𝐷 = (𝑘𝐵𝑇𝑒𝜀0

𝑛𝑒𝑒2)1

2 and is called the Debye length and 𝜀0 is the permittivity of free

space[39]. Debye length is the characteristic plasma dimension on which the charge densities

can exist. From equation 3.2, it can be understood that over a distance 𝜆𝐷 , any plasma

perturbation is reduced to 0.37(1/e) of its initial value. Therefore, the plasma species located

at 2-3 Debye length from the point of perturbation will practically remain unaffected.

38

Assuming 𝑛𝑒 = 1016𝑚−3, and 𝑘𝐵𝑇𝑒 = 2𝑒𝑉 [41],typical value of the Debye length for a cold

low pressure plasma is in the order of 10-4 m [41].

Fig. 3. 3: Plasma potential distribution in a dc glow discharge[39].

Fig. 3. 3 represents the potential distribution in a glow discharge generated by applying a DC

(direct current) continuous voltage at the cathode. The area between the cathode (target) and

anode (chamber walls) is known as the plasma potential (Vp) and has the most positive

potential. In the vicinity of the cathode and anode sheaths are formed, so that the perturbation

of plasma due to the presence of electric is restricted in this region.

3.1.2. Magnetron sputtering

The dc diode base sputtering, discussed in previous section, has several drawbacks. the most

important is that the process requires relatively high sputtering pressure (typically….), which

result in increased scattering of the sputtered species in the gas phase and thus low deposition

rates as well as poor film quality. Further the use of high cathode voltages has a consequence

on the bombardment of anode by fast electrons, which are repelled from the target (cathode)

because of their negative charge. This in return cause substantial heating of the growing film.

39

It is very undesired phenomena which affect the resulting film properties. To avoid such

drawbacks and to be able to maintain the glow discharge at lower sputtering pressures and

voltages, an increase in the ionization of plasma species is essential. This can be achieved by

combining the existing electric field of the cathode with a magnetic field, i.e. by placing a a

set of permanent magnets behind the cathode. In such configuration, the force F experienced

by an electron can be written as

�� = −𝑞(�� + ��×��) (3.3)

where �� is the electric and �� magnetic field vectors, �� is the velocity of the electron and q its

charge. In such scenario, the electron repelled by the negative potential of the cathode will

move helicoidally around the magnetic field lines due to the superposition of the electric field

force �� = 𝑞�� and the Lorentz force 𝐹𝐿 = −𝑞��×��

Fig. 3. 4: the principle of magnetron sputtering. Electron are trapped by Lorentz force in an

inhomogeneous magnetic field and result in increased ionization[42].

The electrons following the helicoidal trajectories are trapped in the vicinity of the target and

result in the increased ionization probability of the gas particles in that region, as shown in

40

Fig. 3. 4. Such configuration in sputtering is known as magnetron sputtering and is the most

popular technique in physical vapor deposition techniques[42,43].

One of the disadvantage of the early developed magnetron sources was that the plasma was

too effectively confined near the target surface due to the magnetic field emerging and

reentering the target in a closed loop type of pattern, also known as balanced magnetic field

configuration[43] (see Fig. 3. 5(a)). In 1986 the issue was resolved by Window and Savvides

[44] by introducing an unbalanced magnetron, Fig. 3. 5(b). In such unbalanced configuration,

stronger outer magnetic ring is used which cannot be compensated by the inner weak

magnetic ring. Such arrangement allows some electrons to escape from the confining

electromagnetic field and create plasma away from the target surface area. This consequently

leads to the better transport of charged particles towards the substrate. Thus enhancing the

ionization in the substrate vicinity allowing for more intense bombardment on the growing

film.

Fig. 3. 5: Schematic illustration of the cross section of (a) balanced and (b) unbalanced magnetron.

41

3.1.3. Reactive magnetron sputtering

Single or multi-component thin films can be deposited by magnetron sputtering only when

the elemental or compound target is used, respectively. In this way, the target has to be made

of the same constituent elements of the material that has to be deposited. This result in the

same stoichiometry of the films as of the target or some time in slight variations, due to

preferential sputtering of the target atoms and/or by the way the sputtered material is

transported to the growing film. Further, this limits the possibility of using magnetron

sputtering where the films composition has to be different than the target. The alternate of

such is the use of reactive gas along with the noble gas (Ar) in the magnetron sputtering

process, and the principle technique is known as reactive magnetron sputtering. Using such

strategy, one can deposit metal oxide, nitride and carbide films by adding O2, N2 or

CH4,/C2H2 in the magnetron sputtering process, respectively. Moreover, it also allows the

thin film scientist to deposit films with various compositions by tuning the flow of the

reactive gas into the deposition chamber.

When a reactive gas is added in the sputtering process, it interacts with the target surface and

with the collecting areas of the chambers, i.e. walls of the chamber and substrate. The

injection of reactive gas has tremendous implications on the sputtering process, as show in

Fig. 3. 6. In the Fig. 3. 6, the reactive gas partial pressure, target voltage and the deposition

rate are plotted during the sputtering of a zirconium target at constant current in a Ar/O2

ambient, as a function of reactive gas (O2) flow[45]. In the beginning, i.e. at very low flow

of O2, the reactive gas partial pressure remains very low (Regime I, Fig. 3. 6(a)) as the O2

flow is increased. This can be attributed to the gettering (the removal of reactive gas atoms

by a getter e.g., Zr) of the reactive gas molecules by the Zr atoms that are freshly deposited

on the substrate surface and the chamber walls [46]. This low O2 partial pressure in regime I

42

results in the very low target coverage (the target surface is not oxidized) with high deposition

rate. As a consequence, the films deposited in this regime contain high metallic content and

is the reason why this regime I is also called metallic mode. The oxygen partial pressure in

the chamber increases abruptly when the oxygen flow rate exceeds the gettering rate, regime

III. This results in the full coverage of the target leading to a ZrO2 compound formation on

the target surface and the decrease in the deposition rate at the substrate. This decrease in

deposition rate is a common feature of reactive magnetron sputtering as the sputtering yield

of formed compound is always lower than the one of the corresponding metal[47]. This could

be explained by the presence of high ionic/covalent bond i.e. high binding energy as

compared to metallic target[48]. This regime III is also called poisoned or compound mode

due to the full coverage/oxidation of the target. In this situation stoichiometric oxide films

are grown. The regime II, lying between the metallic and the poisoned mode, also known as

transition zone, is frequently abrupt and is often accompanied by a hysteresis effect. The

abrupt change in the target coverage and the hysteresis is the result of a nonlinear relationship

between the reactive gas flow and its partial pressure in the chamber as a result of gettering

phenomena. Trying to operate in the regime II with a normal reactive flow control of the

reactive gas is very difficult if not impossible for most reactive systems. It could be also seen

from Fig. 3. 6(b) that such transition from metallic to compound mode has implications on

the target voltage. The increase of target voltage in the case of Zr-O2/Ar indicates the increase

of plasma impedance as the target current was kept constant in this case. Such relation

between the target current-voltage has been studied by Depla et al. [49] for various oxides

and the change in voltage has been attributed to the variation of the secondary electron

emission coefficient. They also found that there is an inverse relationship between the

43

secondary electron emission coefficient and the target voltage, i.e. the voltage drops for

materials where the secondary electron emission increases and vice versa.

As discussed, in the metallic mode films with a very high metallic content are synthesized.

On the other hand, in poisoned mode, only stoichiometric films with very low deposition rate

can be deposited. The only way to deposit films with high deposition rates and with variation

in film chemical composition (under-stoichiometric) is to work inside the transition zone

[30,33,36,47]. The difficulty here, is that this regime is usually unstable and extremely

sensitive to the reactive gas partial pressure. One can overcome this drawback and work

inside the transition zone by using a plasma emission monitoring device or by using a voltage

feedback control unit. In the present thesis work, reactive magnetron sputtering was chosen

for film deposition and the film composition was varied by working inside the transition with

the help of voltage feedback control unit.

Fig. 3. 6: Effect of oxygen flow on the (a) oxygen partial pressure, (b) target voltage and on (c)

deposition rate during the Zr-O2/Ar reactive magnetron sputtering at constant target current [45].

44

3.2. Thin film growth

Once the atoms are sputtered from the target surface, they are transported through the vapor

phase and reach the substrate surface where they condensate and form a film. The

mechanisms involving different processes of thin film formation could be roughly divided

into two stages; i) early stages i.e. the arrival of sputtered atoms onto the substrate surface

(deposition rate, R) and accommodation, their migration along the substrate surface (surface

diffusion) and finally their incorporation into stable clusters (nucleation) and thus 2-

dimensinal (2D) film grow. After initially forming one or more 2D monolayers, further layer

growth becomes energetically unfavorable and 3-dimensional (3D) islands form. The

schematic of these mechanisms is presented in Fig. 3. 7.

Fig. 3. 7: Schematic of processes leading to film growth [50]

45

3.2.1. Early stages of thin film formation

The atoms arriving the substrate from the vapor phase are known as adatoms. When an

adatom lands on the substrate, it loses most of its momentum and kinetic energy through

dissipation into vibrations of the substrate lattice (phonons). In fact, the adatom starts to

interact with the substrate surface when they are at a distance of several Å[41] i.e. adatom

experience an attractive potential of the substrate surface atoms as they approach the

substrate. Thus they gain energy which is in the order of cohesive energy of the substrate

material. Since for the adatom to stick to the surface i.e. not to bounce back to the vapor

phase, it has to reduce its total kinetic energy through dissipative mechanisms to a value lower

than the adsorption energy (Ead). The ratio of the number of deposited adatoms to the number

of impinging adatoms is known as sticking coefficient. Theoretical calculations[51] for

nearly the same masses of adatoms and substrate surface atoms shows that the sticking

coefficient is unity if the kinetic energies are up to 25 times the adsorption energy. Sticking

coefficient are significant lower than unity when high substrate temperature or if atoms

lighter that the substrate atoms are considered. The residence time (𝜏𝑠) of an adatom before

it re-evaporates in to the gas phase is given by equation (3.4)[43].

𝜏𝑠 =1

𝑣𝑜exp (

𝐸𝑎𝑑

𝑘𝐵𝑇) (3.4)

where 𝑣𝑜is the adatom-surface vibrational frequency which depends on the adatom-surface

combination and 𝑘𝐵 is the Boltzmann constant. So an adatom will combine with a surface if

large surface residence time and high sticking coefficients are reached, this criterion is met

only when 𝐸𝑎𝑑 ≫ 𝑘𝐵𝑇.

Once the adatom has landed the substrate surface, the next step is surface diffusion. Surface

diffusion is a very important mechanism for migration of adatoms to the site which allows

46

nucleation and growth. Surface diffusion can be understood as a 2 dimensional random walk

during which an adatom jumps from one potential well to another.

After surface diffusion, next is the nucleation, i.e. the adatoms form aggregates (nuclei) which

grow in size or dissociate. For a relatively low density of nuclei the growth take place via

surface diffusion as shown in Fig. 3. 8. In a first approximation the process depicted in Fig.

3. 8 can be treated thermodynamically and thus the change in free energy on formation of

cap-shaped nuclei with a mean size r can be written as

∆𝐺 = 𝑎3𝑟3∆𝐺𝑉 + 𝑎1𝑟

2𝛾𝑣𝑓 + 𝑎2𝑟2𝛾𝑓𝑠 − 𝑎2𝑟

2𝛾𝑠𝑣 (3.5)

where γ are interfacial energies and v, f and s stands for vapor, film, and substrate,

respectively. The pre-factors, a1, a2 and a3 are geometrical constants which are determined by

the nucleus-substrate wetting angle θ, as shown in Fig. 3. 8. ∆𝐺𝑉 is the change in Gibbs free

energy due to gas-solid transformation, supersaturation in the vapor phase leads to the

negative Gibbs free energy which is the driving force for nucleation.

Fig. 3. 8: Atomistic mechanisms during nucleation[41].

The dependence of ∆𝐺𝑉 on the nuclei mean size r is plotted in Fig. 3. 9. As can be seen from

Fig. 3. 9, there is a critical nuclei size 𝑟∗ on which the ∆𝐺𝑉 depend and thus the nuclei growth.

. The nuclei whose size 𝑟 > 𝑟∗ will grow otherwise, in case of 𝑟 < 𝑟∗ , the nuclei will

dissociate.

47

Fig. 3. 9: Dependence of nuclei free energy on the mean nuclei size r [41]

For the development of a continuous film and its microstructure, the growth of 2-dimensional

nuclei is of paramount importance. Based on interfacial energies, in the case of epitaxial

growth, three growth modes have been established[50], as shown in Fig. 3. 10. The term

epitaxy is used here to describe the growth of a crystalline films on a crystalline substrate.

Depending on the interfacial energies i.e. if 𝛾𝑠𝑣 < 𝛾𝑓𝑠 + 𝛾𝑣𝑓, then the energy balance requires

the minimization of the area covered by the nuclei. This will lead the nuclei to grow in the

form of three dimensional islands, also called (a) Volmer-Weber growth mode (Fig. 3. 10(a)).

In case of ideal homoepitaxy (the case when the film and substrate have perfect lattice match,

i.e. are of same material), 𝛾𝑠𝑣 = 𝛾𝑓𝑠 + 𝛾𝑣𝑓 , will lead to the uninterrupted layer by layer

growth, such growth mode is known as (b) Frank-van-der-Merwe growth mode (Fig. 3.

10(b)). If 𝛾𝑠𝑣 > 𝛾𝑓𝑠 + 𝛾𝑣𝑓, then the area covered by the nuclei should be maximized and will

lead to the formation of one layer at a time, but due to the film-substrate lattice constant

mismatch, strain will develop in the layers. After the growth of few layers, the interfacial

strain energy 𝛾𝑓𝑠 will break the condition for the layer by layer growth mode and will lead to

48

the island growth. Such growth mode is known as (c) Stranski-Krastanov growth mode and

is depicted in Fig. 3. 10(c).

Fig. 3. 10: Schematic of basic growth modes [50].

One has to keep in mind that in sputter deposition, the films are normally grown non-

epitaxially and far from equilibrium. Therefore, the above-mentioned growth mode

classifications are hardly observed. 3-dimensional growth features are more frequently

observed which are predominantly kinetically controlled and result in the production of

polycrystalline films.

3.2.2. 3D thin film growth

The growth stages of a polycrystalline film formation are shown in Fig. 3. 11 [52]. This

evolution of the structure starts by nucleation (Fig. 3.11a) leading to grain growth (Fig.

3.11b), which further leads to the coalescence of the grains either by complete liquid-like

coalescence resulting in the formation of single crystals (Fig. 3. 11c) or by an incomplete

coalescence resulting in polycrystalline islands and channels (Fig. 3. 11d), and finally to a

continuous film (Fig. 3.11e).

49

Fig. 3. 11: Stages of polycrystalline films growth; (a) nucleation, (b) crystal growth, (c) island

coalescence, (d) growth by filling the channels and (e) formation of a continuous film [52].

In such continuous film growth, the same mechanisms are followed; the so-called

fundamental structure forming phenomena. Theses phenomena involve; i) nucleation, ii)

crystal growth and iii) the grain growth[52].

Nucleation involves the same mechanisms as discussed in previous section 3.2.1 i.e. the

condensation of adatoms on the substrate.

Crystal growth involves the incorporation of the deposited material into the condensed phase.

In polycrystalline films two kinds of crystal growth processes can be observed; i) the growth

of discrete crystals which form on the substrate surface, and ii) the growth of crystals on

already formed crystals.

The grain growth in polycrystalline films can also be divided into two categories depending

on the mobility of the grain boundary; i) grain growth by coalescence of the islands, ii) grain

growth by repeated nucleation, also called abnormal grain growth.

50

3.2.3. Microstructure of thin film

Depending on the fundamental structure forming phenomena and the film deposition

conditions, the microstructure of thin film can be described using structure zone model[53].

For physical vapor deposited pure elemental films, structure zone model has been developed

by considering the substrate temperature (Ts) and melting temperature (Tm) of the material

being deposited. Based on the homologous temperature (Ts/Tm), microstructure and growth

evolution of thin film is classified into 3 categories; a) Zone I (0 < Ts/Tm < 0.2), b) Zone T

(0.2 < Ts/Tm < 0.4), c) Zone II (Ts/Tm > 0.4)[53], shown in Fig. 3. 12.

Fig. 3. 12: Structure zone model of pure elemental films as a function of homologous temperature

and film thickness [53].

In Zone I (0 < Ts/Tm < 0.2), the film is composed of very thin fibers of diameters in the range

of 1-10 nm and determined by the nucleation density and statically fluctuations. With

deposition time, the crystalline fibers grow out of the primary nuclei and proceed to the top

of the film while keeping the orientation of the nuclei. However, along the substrate surface,

51

as the Ts/Tm increase, the diameter of the fiber increases. Since in this zone, Ts/Tm <0.2, bulk

and surface diffusion have no remarkable value, the film presents an under-dense structure.

In Zone T (0.2<Ts/Tm <0.4), because of the relatively high temperature, the self-diffusion of

the adatom is remarkable but still the grain boundary mobility is very limited. This result in

inhomogeneous structure along the film thickness, i.e. very fine crystalline at the substrate

but composed of V-shaped grains in the next thickness range.

Zone II (Ts/Tm >0.4), because of the high homologous temperature in this zone, the bulk

diffusion is significant in this regime. This leads to grain boundary mobility not only during

the coalescence but also during the film thickening. The grain boundaries being mobile, the

grain boundary energy minimization also takes place and result in the formation of grain

boundaries perpendicular to the film plane. The films grown in this zone exhibit

homogeneous structure in the growth direction as well as the columnar structure.

Thornton added another parameter in the structure zone model, namely the pressure of the

sputtering gas[54]. This parameter was added to investigate the role of energetic

bombardment on the film microstructure, a basic feature of sputtering process. The model

presented by Thornton predicts that, for films grown at low pressures, the shift of Zone I to

Zone T, or of Zone T to Zone II, to lower homologous temperature (Ts/Tm) due to the

increased energy transfer to the film by the energetic bombardment of plasma species.

As the structure zone model presented above provides the microstructure and growth

evolution of pure elemental thin films grown by magnetron sputtering method. Therefore, the

extension of structure zone model towards the multicomponent and/or multiphase thin films

is needed, especially for the films grown by reactive magnetron sputtering. For the

development of such structure zone model, the determination of the dependence of the

52

fundamental structure forming phenomena on the growth conditions is required. In this

regards, beside the substrate homologous temperature (Ts/Tm), an additional parameter

describing the concentration of contaminant species is considered[53]. Here it should be

noted that the term “contaminant species” does not only describe the impurities which

unintentionally contaminate the films but also the atoms of reactive gas or a second element,

which are intentionally incorporated in the host material lattice for the formation of

compound films. The contaminant species impinging on the growing film can adsorb and

segregate on the faces of the growing crystal or can be dissolved in the crystal lattice[55].

Experimental results show that the contaminant species can either promote or hinder the

operation of fundamental structure forming phenomena[56]. One particular example of such

case is the growth of O containing Al films. Fig. 3. 13 shows the influence of O concentration

on the growing film microstructure. In general oxygen has a low solubility in Al and

segregates on the surface and at the grain boundaries. This lead to the 2-dimensional layer

formation at these sites and hinders the surface diffusion and grain boundary mobility. At

low O/Al flux ratio (Jo/JAl ~ 10-3), O accumulates at the grain boundary and hinders the grain

boundary mobility. This result in zone II films with a poor texture compared to pure Al thin

films (Fig. 3. 13 a & b). At higher O concentration (Jo/JAl ~ 10-2), coarsening during the

coalescence is partially suppressed which lead to the formation of randomly oriented grains

i.e. zone T films (Fig. 3. 13 c). On further increase of O concentration (Jo/JAl ~ 0.1-1), an

oxide layer form on all islands of all possible orientation, leading to the periodic interruption

of grain growth and the start of secondary nucleation. This result in nano-size grains (Fig. 3.

13d) and is also referred as to zone III. This zone III is also predicted by structure zone models

when high deposition temperatures close to the melting point are used[53]. Finally, when the

53

O atoms are in majority as compared to Al atoms i.e. Jo/JAl > 1-5, then the film consists

mainly of Aluminum oxide which is an amorphous phase at room temperature (Fig. 3. 13e).

Fig. 3. 13: Influence of oxygen concentration on Al thin film microstructure[56].

3.3. Influence of energetic species on thin film properties during

sputtering

An interesting feature of plasma-based sputtering is the bombardment of the growing film by

energetic plasma species. As discussed previously, plasma species can be ions, electron,

photons and as well as neutrals [41]. The interaction of these plasma species especially of

ions with the growing film is determined by their kinetic energy and flux. [38]. Such

54

interaction of bombarding ions with the growing film depending on their energy is shown in

Fig. 3. 14. Of course the exact energy values depend on the growth conditions. It could be

seen in Fig. 3. 14, energy range span from few eVs to 1000 eV. The lower limit of the scale

i.e. below 0.1 eV describes the thermalized species which do not affect the growing film. The

activation of surface process starts from 0.1 eV, where physisorption of impinging species is

possible. An order of magnitude higher, i.e. at 1 eV, chemisorption processes start where the

chemical reactions are involved and last up to 10 eV. The surface characteristics of growing

film surface starts to get affected when these energies reach tenths of electron volts. Finally,

when the energies reach 100th of eV, the impinging atoms start to influence the bulk of the

growing film. The upper limit of the energy is defined where the kinetic energy and

momentum of the impinging species is so high that the deposited material is again vaporized

(re-sputtering). This effect starts about 1000 eV, depending on the material.

Fig. 3. 14: Schematic of effect of energetic bombardment at different energy levels, after [38].

In the energy range of chemisorption, the depositing ions are largely ineffective to influence

the film physical properties. On the other hand, the energy range of surface and bulk effects

are the most important as they provide additional means to tailor the film properties. For

55

example, in the energy range of surface effects (tenth of eVs), depositing ions are only able

to influence the atoms which are in their vicinity of impinging position. This will not cause

any damage but will support the annealing of defects frozen in as a result of limited adatom

mobility. On the other hand, in the energy range of bulk effects i.e. 100th of eVs, the

impinging ions get implanted in the subsurface of the growing film and cause subsurface

effects. These implantation leads to the lattice defects e.g., displacement of atoms from their

lattice site and creation of vacancy. These phenomena lead to the generation of residual

compressive stresses, which are discussed in next section. Sometime these implantations can

also cause the collision cascade, which result in the re-sputtering of the deposited material.

Moreover, these high energetic impinging ions can also penetrate deep in the growing film,

known as channeling mechanism. In this case, the ballistic damage is significantly reduced

and the energy of the impinging ions is minimized by electronic excitations. Not always all

the high energetic bombarding ions are implanted in the growing film, sometime they are

also backscattered. Finally, the interaction of bombarding ions with growing film surface can

also cause electron and photo emission.

The influence of bombarding ions has also shown implications on the crystal structure and

microstructure of the zirconia films prepared by reactive magnetron sputtering [28,31,32,57],

as already discussed in section 2.2.2. It has been shown that at low sputtering pressure or at

less target-substrate distances or at high target powers (i.e. at high energy) the deposited

zirconia films exhibit dense microstructure, which turn porous on the increase of sputtering

pressure. Increasing the target-substrate distances or at lowering the target power (i.e.

decreasing the ion density) have an identical effect on the zirconia film microstructure. One

also has to note this the data presented here deal with the energy brought to the growing film

by the ions only. But in reality the total energy brought to the growing film is a combination

56

of heat radiations (photons) towards the substrate, power transferred by electrons and ions,

neutral depositing atoms and as well as of background gas [58–61]. Taking into account these

contribution, the total energy flux per deposited metal atom can reach several keV [61].

Further, it has also been observed that, with the increase in impinging atoms energies, films

exhibit an increase in compressive stresses. Not only this, with the change in imping ion

energy a change in film crystal structure is also observed i.e. pure monoclinic or mixed

tetragonal and monoclinic or pure cubic can be obtained.

3.4. Formation and evolution of stresses in thin films

Films grown by physical vapor deposition method experience a continuous bombardment of

energetic plasma species, as discussed in the previous section. This leads to the enhanced

grain boundary mobility as well as in the formation of defects in the growing films.

Fig. 3. 15: Schematic of (a) tensile and (b) compressive stresses in the deposited film [62].

As a result, mechanical stresses are generated in the thin films and is called intrinsic stresses.

Such generated intrinsic stresses affect the films crystallinity by causing lattice distortion and

57

sometime these stresses reach so high value that they cause the phase transformation [28].

Moreover, this also affect the mechanical, thermal and electrical properties of the deposited

films. In general film intrinsic stresses can be divided into two parts; a) tensile stress and b)

compressive stress, shown in Fig. 3. 15. In polycrystalline thin films, in the post coalescence

growth stages, grains interact in order to minimize their energy and thus close any existing

gap. This leads the grain boundaries to shrinkage, which in turn causes tensile stresses [63].

Magnitude of such tensile stresses can be calculated using the equation (3.6)[41].

𝜎𝐺𝐵 =𝐸𝑓𝐷

𝐿𝑔(1 − 𝑣𝑓)(3.6)

Here 𝐸𝑓 is the elastic modulus, 𝑣𝑓 is the Poisson’s ratio of the film, 𝐿𝑔 the grain size and D

is the grain boundary relaxation distance. Similar to grain boundary shrinkage, shrinkage of

film volume due to the recrystallization, phase transformation and defect annihilation cause

tensile stresses[64]. Contrary to the tensile stresses, compressive stresses are generated due

to the expansion of the crystal lattice. In PVD, the lattice expansion is caused by the

implantation of energetic plasma species, as discussed earlier. However, compressive stress

relaxation beside the compressive stress generation at high energies is also observed

frequently[65].

Fig. 3. 16: Stress dependence on the energy of bombarding species [66].

58

The rationale behind it is, when the energy of the bombarding specie is too high, this cause

the intense local heating of the film (known as thermal spike). This provides enough energy

to the implanted and displaced atoms to move from their metastable position to film surface

(out-diffusion). Since these tensile and compressive stresses in the film are dependent on the

adatom energy, three energy regimes related to the residual stresses can be considered as

shown in Fig. 3. 16[66]; (i) low energy regime (0.1-1 eV), in this regime tensile stresses due

to the grain boundary shrinkage are dominant. (ii) moderate energy regime (1-20 eV), in this

energy regime compressive stresses are dominant due to the plasma species implantation and

atom displacement. (iii) high energy regime (>25 eV), in this energy regime stresses

relaxation occur due to the thermal spikes.

The intrinsic stresses in the film deposited on a substrate lead to the change in the curvature

of the substrate (δ), as shown in Fig. 3. 15. Using Stoney’s equation and change in substrate

curvature and film, substrate mechanical properties, one can calculate the values of those

intrinsic stresses generated in the film [67].

Beside intrinsic stresses caused by bombardment of energetic species, stresses at the film-

substrate interface can also be generated, when the film is deposited at a different temperature

than the surrounding temperature. For example, deposition of film at high temperatures. Such

stresses are called thermal stresses and originate because of the difference in film and

substrate coefficient of thermal expansion. Thermal stresses in a film can be calculated using

equation (3.7) [68]

𝜎𝑡ℎ =𝐸𝑓

1 − 𝑣𝑓(𝑎𝑓 − 𝑎𝑠)(𝑇𝑠 − 𝑇𝑎) (3.7)

59

here 𝐸𝑓 is the elastic modulus, Ts is the substrate temperature during deposition, Ta is the

temperature during measurement, 𝑣𝑓 is the Poisson ratio, 𝑎𝑠 and 𝑎𝑓 are thermal expansion

coefficient of substrate and film respectively. Thermal stresses can be compressive or tensile

depending on the sign of 𝜎𝑡ℎ. If 𝜎𝑡ℎis positive, then the stresses are tensile and if 𝜎𝑡ℎnegative,

then the stresses are compressive.

Beside stress generated by the film depositing species or by the difference in film-substrate

thermal expansion coefficient, evolution of film stresses has also been observed as a function

of film thickness [67,69]. In this regard Doerner et al. has discussed two hypothetical cases,

i) the stress increase continuously as the film grows i.e. a linear relation between the stress

and the film thickness and ii) the stress grow as function of film thickness up-to a certain

value and then stabilize. Nouveau et al. [69] have observed a similar case to (ii) for CrN films

deposited by magnetron sputtering as a function of film thickness. and observed three growth

regimes: Regime A) (up to 150nm) where the stress keeps increasing, B) (150-500nm) the film

changes its structure to release stress, C) (above 500nm) the stress value is stable and relaxation is

induced by the growth of less dense planes.

60

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4. Zirconia thin film deposition and

characterization

Today several theoretical and experimental methods exist study the growth of thin films (as

discussed in chapter 3) and their properties. In this thesis work, Density Functional Theory

(DFT) calculations are used to better understand the influence of oxygen vacancy

incorporation on zirconia phase formation. This theoretical investigation is then compared to

the properties of oxygen vacancy doped zirconia thin films synthesized by cold plasma-based

dc reactive magnetron sputtering (dc-RMS). The choice of reactive magnetron sputtering is

not only based on its simplicity of use but also because it allows controlling the film elemental

composition e.g. with the help of a voltage feedback control unit. To characterize the

deposited zirconia thin films, several analysis techniques are used. Obviously, the choice of

the analysis methods depends on the properties one wishes to investigate, the resources

available, and on the precision required.

In this chapter, the technical details related to the DFT calculations and thin film deposition

are summarized. Moreover, a description of the thin film characterization methods used to

address the film phase constitution and microstructure as well as the ionic conductivity and

photoluminescence properties is also provided.

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4.1 Modelling and computational details

Theoretical calculations are performed at the DFT level using the 3.2 version of the SIESTA

code [1] including periodic boundary conditions. The calculations are made in the

Generalized Gradient Approximation (GGA) using the Perdew-Burke-Ernzerhof (PBE)

functional [2]. Trouillier-Martins pseudopotentials were used to simulate the nucleus and

core electrons. We used 5s2 4d2 as valence configuration for Zr with 3.04, 3.19 and 2.68 Bohr

for s, p, and d channels, respectively. For oxygen, we have used 1.14 Bohr for all channels.

The atomic basis set is described using a double zeta basis plus polarization orbitals (DZP).

Supercells of cubic, tetragonal, and monoclinic polymorphs of ZrO2 are built using 96 atoms

i.e. 32 Zr and 64 O by a duplication of the experimental unit cell in the three directions. The

final lattice vectors of the three polymorphs are: a = 10.14 Å for cubic; a = b = 10.10 Å, c =

10.36 Å for tetragonal; a = 10.29 Å, b = 10.42 Å, c = 10.62 Å, β = 99.23° for monoclinic.

Oxygen vacancies are inserted in the zirconia lattice by removing O atoms from the pristine

ZrO2 cells. Three configurations were tested. The vacancies are distributed randomly, apart

from each other, and as clusters, to assess the influence of O vacancies on the phase

constitution. In case of random vacancy generation, O atoms were first assigned numbers and

then from those total assigned numbers some number were randomly picked by the software

and O atoms corresponding to those numbers were removed. In case of apart O vacancies

were created manually by removing O atoms apart from each other. In case of cluster 2-4

vacancy were clustered together.

The optimization of the structure and the calculation of the electronic structure are performed

at 0 K using a mesh cut-off of 190 Ry and a Monhkorst-Pack grid of (2x2x2) k-points. The

atomic positions were relaxed until the forces on the atom were less than 0.04 eV/Å, while

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the lattice vectors of the unit cell were fixed during the calculation. Moreover, the density of

states (DOS) of cubic structure were also extracted from DFT calculations data, for various

O vacancy concentrations.

4.2 Thin film deposition and process monitoring

To assess experimentally how the incorporation of O vacancies influences the zirconia phase

formation/stabilization, experiments were designed using a conventional dc reactive

magnetron sputtering (dc-RMS) setup where a Zr target is sputtered in a reactive atmosphere

containing argon and oxygen. The sputtering setup is built in such a way that it not only

helped to incorporate O vacancies in the zirconia lattice as the film grows, but also to vary

their concentration.

For this purpose, a voltage feedback control unit6 [3–6] (Speedflo mini from Gencoa, UK)

was used. The voltage feedback control unit is an auxiliary device which provides a rapid

control over the oxygen partial pressure using the target voltage as a feedback signal. This

way, the reactive sputtering process was allowed to work inside the so-called metal-to-

compound transition zone [7] of the sputtered zirconium target. The transition zone is an

experimental working window in between the metallic and oxidized modes of the sputtered

target where under-stoichiometric films are grown, as shown in Fig. 4. 1. Using this well-

optimized synthesis setup, 100 ± 10 nm thick films were deposited on Si (100) single crystal

substrates at 50, 65, 75, 80, and at 100% signal set point values. The 0 % represents the

metallic mode set point while 100 % represents the peak of the transition zone (See Fig 4.1).

Film were also grown in the oxidized mode. In order to reach a precise control over the sputter

target and therefore of the film-chemistry, O2 gas was injected at the target during the film

6 Discussed in details in section 4.2.1

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growth and the length of the tubing system, between the O2 mass flow and the vacuum vessel,

was shortened as much as possible. It should be noted that the films used to perform the

elemental analysis, 18O2 (purity 97.1 %) was used instead of conventional 16O2. The Ar gas

was injected away from the target by using a conventional mass flow controller. A schematic

representation of the zirconia thin film deposition process is shown in Fig. 4. 2 and a summary

of deposition conditions is given in Table 4. 1.

Fig. 4. 1: Target voltage curve of Zr target as a function of O2 flow, shows the transition zone and

working points inside the transition zone as well as in the poisoned zone where the ZrO2-x and ZrO2

films were deposited, respectively.

Table 4. 1: Experimental details of the film deposition by reactive magnetron sputtering.

Base pressure <2×10-6 mTorr

Working pressure 5, 10, 20 mTorr

Target to substrate distance 6.5 cm

Discharge current (I) 0.2, 0.3, 0.4 A

Reactive gas 18O2 or O2

Sputtering gas Ar

Substrate Si (100); 525 ±20 μm

thick

Film thickness 100±10 nm

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Every deposition run was carried out in a high vacuum chamber (length 78 cm, diameter 42

cm). The base pressure < 2×10-6 mTorr (4×10-4 Pa) was reached with help of Turbo pump

backed by primary pump (pumping speed 12 m3/h). During the depositions, the working

pressure was kept constant at e.g. 10 mTorr (1.33 Pa) in each case using a throttle valve. To

introduce sample in the deposition chamber, a load lock system was used and before

depositing each film, the target was first sputtered clean. During the depositions, non-

intentionally heated substrates were placed at a distance of 6.5 cm from a 5 cm in diameter

purity (99.97%) Zr target. The DC current applied to the sputter target (i.e., the cathode of

the system) was ranging from 200 mA up to 400 mA. The power supply was an Advanced

Energy MDX 500 dc power supply equipped with an arc suppressor (Sparkle from Advanced

Energy).

Fig. 4. 2: Schematic of Zirconia thin film deposition used in this work. The target voltage is read by

the feedback control unit (right hand side) and allows controlling the oxygen flow during film

growth.

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4.2.1 Voltage feedback control unit

The voltage feedback control unit (Speedflo mini from Gencoa, UK) is a device which

actively monitors reactive magnetron sputtering processes and controls them by continuously

monitoring the target voltage or other variables. The principle of the voltage feedback control

unit is similar to a basic automatic control system which is described by a block diagrams in

Fig. 4. 3. When the user of the automatic control system inputs a reference value (r) which

has to be achieved (e.g. the temperature of a room), then the controller having an algorithm

will send some manipulated variable (m1) to the control elements (for example a heater). The

final control elements will then give a new manipulated variable (m2) (for example heat flow).

That variable will be (eventually) combined with other loads (l) from the system (e.g. sun or

human heat) which are not manipulated, ultimately giving an overall manipulated variable of

the reaction m3. By having a way of monitoring what the reaction was (for example measuring

the temperature in the room) one can get what is called the controlled variable value c. This

value c is fed back to the controller which compares it with the reference value (r). Based on

the error ‘e’, controller will act consequently in the loop control.

Fig. 4. 3: The basic automatic control system [3].

In voltage feedback control units, the target voltage is monitored continuously and is used as

input. Input is processed by the embedded software and an actuator output signal is produced,

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based on certain algorithm parameters. The actuator output signal is sent to the mass flow

controller which responds according to the signal and regulates the reactive gas flow to

maintain the set point value condition. For a reliable voltage feedback control, the system

requires effective sensor information. For this purpose, voltage feedback system generates a

series of commands to the actuator via a processing algorithm which should be appropriate

in relevance, magnitude and speed of change.

The speed of data transfer between the voltage feedback control used here and the PC

(Windows) is called Update Speed. This parameter can be varied between a minimum of 0.5

to 5 sec depending on the working situation. In this thesis work, the highest update speed 0.5

seconds is used. There is also the Sampling/Actuation rate of the voltage feedback unit that

should not be confused with the Update speed above. The sampling/actuation is related to the

Input/Output communications with sensors and actuators that are in the millisecond range.

The voltage feedback controller used in this thesis work can handle large numbers of data

sampling with typical rates between 1-10 ms depending on the number of mass flow

controllers connected and the type of action on each process. In the thesis work, only one

mass flow was connected to the process and the sampling response type during the process

was selected as AUTO, allowing the voltage feedback unit to operate as fast as the processing

time allows (around 1 ms or less).

Beside using voltage feedback control unit to work inside the transition zone, there are other

options available as well e.g., plasma emission monitoring (PEM) and lambda probe. In case

of PEM, the intensity of an emission line of the element of interest present in the plasma is

monitored and used to work inside the transition zone. To use PEM, one needs to collect the

70

plasma emission efficiently and in a robust way as any variation in the line of sight or

deposition of sputtered material on the collecting window can lead to misleading conditions.

The lambda probe or lambda sensor gives a measure of the partial pressure of oxygen with

respect to the atmosphere. Lambda probe is directly installed in the process chamber. It is

based on the Nerst differential voltage potential created when the process side of the probe is

exposed to a different partial pressure of oxygen than the atmospheric side of the probe. The

probe needs O- diffusion in order to generate the potential and achieves this via a platinum

catalyst and a zirconia membrane. Since the zirconia and platinum require high temperatures

in order to be operational, therefore the sensor is also fed with an independent electrical

power. The choice of Voltage feedback unit in this thesis work is based on its simplicity to

use and the geometrical restrictions of the chamber.

4.3 Film characterization tools

To characterize the deposited zirconia thin films, various characterization techniques are used

in this thesis work and are briefly presented in this section.

4.3.1 X-ray diffraction (XRD)

X-ray diffraction (XRD) [8] is a powerful technique vastly used in material science division

to identify the crystalline phases present in materials and to measure their structural properties

(grain size, phase composition, stress). The sample to be analyzed is irradiated by x-rays

(having wavelength in the range of interatomic spacing in the crystals). The X-rays interact

with the matter, get diffracted, and are collected by a detector. Diffraction of x-rays by

periodic arrays of atoms in the crystals result in constructive and destructive interferences

which give rise to the diffraction pattern. This phenomenon is described by Bragg’s Law [9]

2𝑑𝑆𝑖𝑛𝜃 = 𝑛𝜆 (4.1)

71

where, d is the lattice spacing, 𝜃 is the scattering angle, 𝜆 is the wavelength of x-rays and n

is an integer number. During XRD measurements, the sample is scanned for a wide range of

angles and the diffracted intensities are recorded as a function of the latter. Such recorded x-

ray pattern is called diffractogram. Recorded diffractogram of the crystalline material is

viewed as a fingerprint of the crystal structure and is compared to the diffractograms in the

database to identify the material crystal structure. Beside identifying the crystal structure, one

can also extract the information regarding the grain size, crystal orientation, phase

composition and stresses present in the material by using the peak position, peak intensity,

peak width, and the change in peak position

4.3.1.1 Bragg-Brentano (θ-2θ) mode

The diffractograms are typically recorded by using the measurement setup in the so-called

Bragg-Brentano (θ-2θ) [8] geometry, where θ is the incidence angle of the x-rays measured

with respect to the surface of the sample and 2θ is the diffracted angle. In this configuration,

the source of the x-rays beam and the detector are scanned synchronously such that the

incidence angle and the diffracted angle remain the same throughout the scanning. Therefore,

in such a geometry only the crystal planes which are parallel to the sample surface are probed.

A schematic of such Bragg-Brentano (θ-2θ) geometry is shown in Fig. 4. 4(a).

Fig. 4. 4: Schematic of (a) Bragg-Brentano (θ-2θ) and (b) GIXRD geometry.

72

4.3.1.2 Grazing incidence XRD (GIXRD) mode

Normally in Bragg-Brentano (θ-2θ) mode, during the x-ray diffraction of thin films

(thickness in the range of tens of nm), the intensity of the diffracted peaks is very low as well

as the diffractograms are dominated by diffractions from the substrate due to the larger

penetration depths of x-rays (of the order of µm). This is particularly the case when highly

ordered substrates such as Si single crystals are used. In order to avoid such kind of issues,

the method of grazing incidence x-ray diffraction (GIXRD) [8] is used. During GIXRD, the

x-rays enter the sample at very small fixed incidence angles (few degrees or less) of incidence

thereby increasing the path travelled by the x-rays significantly (see the schematic shown in

Fig. 4. 4(b)). Since the incidence angle (𝜔 in Fig 4.4.) is fixed in GIXRD mode, therefore,

the diffractograms are obtained by varying the detector position. This allows probing crystals

whatever their orientation in the film.

In the present thesis work, crystallinity of the deposited films was analyzed by BB method

and GIXRD using PANalytical Empyrean with a Cu-Kα radiation (λ = 1.54 Å) source. The

diffractograms were recorded with a step size of 0.07° using an incidence angle of 0.5° at 40

mA, 45 kV of generator settings. The resulting diffractograms were compared to ICDD PDF

cards representing patterns of three polymorphs of ZrO2: monoclinic (PDF# 37-1484),

tetragonal (PDF# 81-1544) and cubic (PDF# 49-1642) in order to know the crystal structure

of the deposited films.

4.3.2 Transmission electron microscopy (TEM)

Transmission electron microscopy (TEM) [8] is a powerful technique for examining the

crystal structure and the microstructure of solid materials e.g., metals, ceramics,

semiconductors, polymers, and composites. In TEM, a focused electron beam with energy

(100-300 kV) is incident on a thin (typically less than 200 nm) sample. The signal in TEM is

73

obtained from both un-deflected and deflected electrons that penetrate the sample. A series

of magnetic lenses at and below the sample position deliver the signal to a detector, usually

a fluorescent screen, a film plate, or a video camera. Accompanying this signal transmission

is a magnification of the spatial information in the signal by as little as 50 times to as much

as a factor of 106. One considering factor regarding the TEM lenses are the diaphragms or

apertures of these lenses employed at certain positions in TEM. The purpose of these

apertures is to filter either the source or the transmitted signal. The most important diaphragm

in TEM is called the objective aperture, lying at the back focal plane of the objective lens. In

this plane the scattered electron waves recombine to form a diffraction pattern. The use of a

small objective aperture lenses while operating in the image mode, blocks all diffracted

beams, and serve to enhance the image contrast significantly. The use of a large objective

aperture allows the passage of many diffracted beams thus enhanced diffraction pattern, is

the modus operandi for the technique and is referred to high-resolution transmission electron

microscopy (HRTEM).

In the present thesis work (HR)TEM is used to study the crystal structure of zirconia films at

various positions from the substrate-film interface.

4.3.3 Secondary electron microscopy (SEM)

Scanning electron microscopy (SEM) [8] is one of the most widely employed techniques to

the study the surface topography in three dimensions of a material. In SEM, a beam of

electrons, like in TEM, with energy ranging from few keV to 50 keV is emitted from a W-

filament or LaB6-crystal.The beam is focused on the surface of the sample with the help of

magnetic lenses. As the electrons impinge on the surface, interactions occur. As a result of

this, the emanating secondary and backscattered electrons are collected by suitable detectors

74

and are used to extract the microstructure and surface topography information in the form of

images of the sample.

In the present thesis work SEM is used to study the microstructure of the deposited samples

by having their cross-sectional images.

4.3.4 Chemical composition of zirconia thin films

The chemical composition of zirconia thin films i.e. concentration of Zr and O in this thesis

work was realized by combining Rutherford backscattering spectrometry (RBS) [8] and

Nuclear reaction analysis (NRA) [8].

4.3.4.1 Rutherford backscattering spectroscopy (RBS)

Rutherford Backscattering Spectroscopy (RBS) is one of the commonly used nondestructive

technique for quantitative depth-profiling of thin films. In RBS, sample is bombarded with a

beam of high-energy (Eo) particles (MeV range) with mass M1, which undergoes an elastic

collision with the sample stationary atoms to be investigated (having mass M2). On collision,

the energy of backscattered M1 particle is detected at a given angle θ. By using laws of

conservation of energy and momentum, the mass of the target particles M2 is calculated.

Moreover, since the probability of scattering in a certain angle is known by the so-called

Rutherford cross section, this makes it possible to estimate the abundance of M2 by counting

the yield of scattered particles M1 in a certain solid angle, covered by the detector.

Results obtained by RBS are insensitive to sample matrix and typically do not require the use

of standards, which makes RBS the analysis of choice for depth profiling of major

constituents in thin films. Detection limits range from a few parts per million (ppm) for heavy

elements to a few percent for light elements.

75

In the thesis work, for elemental analysis, zirconia thin films were deposited on graphite foil.

The concentration of Zr in the samples was probed by RBS. The incident energy of the alpha

particles was 2 MeV, and the beam impinged the sample surface at normal incidence.

Backscattered particles were collected at 165° in a passivated implanted planar silicon (PIPS)

detector. Spectra were analyzed with the SIMNRA software assuming Rutherford

backscattering cross-section.

4.3.4.2 Nuclear reaction analysis

While RBS is more suitable for heavier elements, Nuclear reaction analysis (NRA) is used to

determine the concentration and depth distribution of lighter elements in solids but like RBS,

NRA is also a nondestructive technique. NRA is also isotope specific, making it the ideal tool

for isotopic tracer experiments. This characteristic also makes NRA less vulnerable than

some other ion scattering methods to interference effects that can overwhelm signals from

low abundance elements. In NRA, an ion beam with an energy ranging from a few hundred

keV to several MeV is produced in an accelerator and bombards the sample. Nuclear

reactions with low-Z nuclei in the sample are induced by this ion beam. Products of these

reactions (typically p, d, t, He, α particles, and γ rays) are detected, hence producing a

spectrum of particle yield versus energy. That accumulated spectrum is then compared to a

body of data accumulated through research in low-energy nuclear physics to determine

concentrations and distributions of specific elements or isotopes in the material under

investigation. Depending on the reaction type, NRA can be divided in to two different

categories: The Resonant Nuclear Reaction Analysis (RNRA) and the non-resonant Nuclear

Reaction Analysis. RNRA uses beam energies near narrow isolated resonances of relevant

nuclear reactions to determine the depth distribution of elements in a sample. On the other

hand, when reaction cross sections are sufficiently large over an extended energy range, the

76

entire depth profile may be obtained using a single incident beam energy. This is known as

non-resonant NRA.

In this thesis work for elemental analysis, zirconia thin films were deposited on graphite foil.

Specific 18O depth profiles were determined using the resonant nuclear reaction 18O(p,α)15N

at 151 keV. In this case, in order to allow NRA of the zirconia films, isotopic 18O was

introduced in the vacuum chamber instead of conventional high-purity oxygen. The latter is

being 16O. This way, only the oxygen introduced in the film during deposition is probed and

oxygen pollution from venting is discarded. Samples were tilted at 30° with respect to the

incident beam and the alpha particles were collected in large area passivated implanted planar

silicon (PIPS) detectors facing the sample surface and parallel to it. The incident energy was

varied from 145 to 200 keV. Depth profiles were deconvoluted in order to take into account

the energy straggling of the beam and then quantified with the help of a Si18O2 standard

produced by thermal oxidation in a pure 18O atmosphere. These RBS and NRA measurements

were performed in the Laboratoire d’Analyse par Réaction Nucléaire (LARN, Pr. S. Lucas)

at the University of Namur.

4.3.5 Electrochemical Impedance spectroscopy (EIS)

Electrochemical Impedance spectroscopy (EIS) [10] is a well-known method for

characterizing many of the electrical properties of materials with electronically conducting

electrodes. For example, it can be used to investigate the dynamics of a bound or mobile

charge in the bulk or interfacial regions of a solid or liquid material (e.g., electrolyte,

semiconducting, mixed electronic–ionic). EIS involves the measurement of impedance by

applying a single-frequency voltage to the cell (i.e. the electrolyte equipped with the

electrodes) and measuring the phase shift and amplitude, or real and imaginary parts, of the

resulting current at that frequency using either an analog circuit or fast Fourier transform

77

(FFT) analysis of the response. Later the measured impedance of a cell is compared to an

equivalent RC model circuit, leading to the electrolyte resistance i.e., impedance R. Using

this measured value of resistance and the geometry of the cell, one can calculate the ionic

conductivity using equation 4. 2.

𝜎 = 1

𝑅

𝑎

𝑏. 𝑑(4.2)

Where, 𝜎, is the ionic conductivity of the electrolyte, a, is the distance between the two

electrodes, b, is the electrode width normal to the current flow and d is the film thickness, as

shown in Fig. 4. 5.

Fig. 4. 5: Schematic of film on a substrate of which impedance is measured.

In the present thesis work, EIS was performed on the deposited samples to measure their

lateral ionic conductivity in air. Measurements were carried out in a two-electrode

configuration geometry using silver paste to ensure the connectivity onto the film surface.

EIS measurements were performed in the 475 – 725 °C temperature range with 25 °C steps.

During the measurement, the frequency was varied from 42 Hz to 1 MHz, with a 0.2 V

alternating voltage signal.

4.3.6 Photoluminescence (PL)

Luminescence is referred to the emission of light by a material through any process other

than blackbody radiation. The term Photoluminescence (PL) [8] narrows this down to any

78

emission of light that results from the optical stimulation of the material. In our everyday life,

PL finds its uses from food to materials research area. For example, many inorganic materials

including, semiconductors, crystalline ceramics, and glasses are studied by PL to probe their

optical properties and to have a look at impurities and defects.

In PL, when a material is irradiated it gains energy by absorbing the photons at some

wavelength and an electron is excited from a low to a higher energy level. Such a process can

be described as a transition from the ground state to an excited state of an atom or molecule,

or from the valence band to the conduction band of a semiconductor crystal (electron-hole

pair creation). After excitation, the system may undergo a non-radiative internal relaxation

involving interaction with crystalline, and then the excited electron moves to a more stable

excited level, such as the bottom of the conduction band. After a system-dependent

characteristic lifetime in the excited state (which can last from picoseconds to many seconds)

the electron returns to the ground state by emitting the excess energy in the form of light

(known as radiative transition). This emitted light is detected as photoluminescence, and the

spectral dependence of its intensity is analyzed to deduce the information about the properties

of the energy states of the material. The time dependent emission can also be monitored to

get informations about energy levels coupling and lifetimes. The light used in PL excitation

and emission usually fall in the range of 0.6-6 eV (roughly 2000-200 nm). Many electronic

transitions of interest lie in this range, and efficient sources and detectors for these

wavelengths are available easily.

In this thesis work, photoluminescence (PL) measurements were carried out at room

temperature in order to investigate optical properties of OVSZ samples. For this purpose, a

Xenon lamp was used as excitation source and an edge filter was used as to cut the excitation

79

wavelength above 325 nm. During the characterization, excitation wavelengths used are 270

nm (4.6 eV), 280 nm (4.4 eV), 290 nm (4.3 eV), 300 nm (4.1 eV), and 310 nm (4 eV).

Detection was made possible using a CCD camera and, in order to compare the PL spectra,

PL intensities were normalized to the excitation power.

4.3.7 Heat flux microsensor

In order to investigate the role of energetics on the growing film, a Heat flux microsensor

(HFM) developed by the GREMI research group from the University of Orléans (Dr. A.-L.

Thoomann [11–15]) was used. The probe was mounted at the substrate position, 6.5 cm above

the sputter target. The heat flux sensor converts heat into an electrical signal [11–15] and is

recorded as an energy density (W/m2) delivered to the substrate – film system during

deposition. Such measured energy actually represents the “global energy” delivered during

the deposition process, it detects simultaneously the contributions of i.e. Ar+ and other plasma

ions bombarding the film, fast particles such as O- or secondary electrons emitted from the

target, electrons, photons, IR photons emitted by the hot sputtered target [13], condensing

sputtered atoms, the enthalpy of formation of metal-oxygen bonds. In this study, the energy

density recorded by the HFM was normalized with respect to the deposition rate in order to

allow the comparison between various deposition conditions.

The heat flux instrument used is built around a commercial sensor from Vattel (Vattel-HFM-

7) and is composed of a thermopile for the energy flux measurements and a resistance

temperature detector (Pt100) for the temperature control. The thermopile used contained

1600 thermocouple junctions per cm. The small diagnostic active area (approx. 17mm) and

the ms time resolution provided a very sensitive probe for energy flux measurements. The

device was calibrated according to NIST protocols using a cylindrical black body. During the

measurement, the temperature of the HFM was maintained around 5 °C using a proper

80

cooling system in order to avoid the heating of the sensor as this can lead to energy loss by

IR emission. Before starting the measurement, a 100 μm thick copper foil used as substrate

was pasted on the HFM sensitive surface using a vacuum-compatible thermal paste (JELT-

6017-COMPOUND SILICONE 20G). The presence of the copper foil only influences the

response time of the system. In order to avoid the growth of thicker film on the copper foil,

the latter was changed periodically.

81

References [1] J.M. Soler, E. Artacho, J.D. Gale, A. Garcìa, J. Junquera, P. Ordejòn, D. Sànchez-Portal, The SIESTA Method

For Ab Initio Order-N Materials Simulation, J. Phys. Condens. Matter. 14 (2002) 2745–2779. doi:10.1088/0953-

8984/14/11/302.

[2] J. Perdew, K. Burke, Y. Wang, Generalized gradient approximation for the exchange-correlation hole of a many-

electron system, Phys. Rev. B. 54 (1996) 16533–16539. doi:10.1103/PhysRevB.54.16533.

[3] V. Bellido-González, B. Daniel, J. Counsell, D. Monaghan, Reactive gas control of non-stable plasma conditions,

Thin Solid Films. 502 (2006) 34–39. doi:10.1016/j.tsf.2005.07.230.

[4] W.D. Sproul, D.J. Christie, D.C. Carter, Control of reactive sputtering processes, Thin Solid Films. 491 (2005) 1–

17. doi:10.1016/j.tsf.2005.05.022.

[5] D. Depla, G. Buyle, J. Haemers, R. De Gryse, Discharge voltage measurements during magnetron sputtering,

Surf. Coatings Technol. 200 (2006) 4329–4338. doi:10.1016/j.surfcoat.2005.02.166.

[6] I. Safi, Recent aspects concerning DC reactive magnetron sputtering of thin films: a review, Surf. Coatings

Technol. 127 (2000) 203–218. doi:10.1016/S0257-8972(00)00566-1.

[7] J. Musil, P. Baroch, J. Vlček, K.H. Nam, J.G. Han, Reactive magnetron sputtering of thin films: Present status

and trends, Thin Solid Films. 475 (2005) 208–218. doi:10.1016/j.tsf.2004.07.041.

[8] C. Richard Brundle, C.A. Evans, S. Wilson, eds., Encylopedia of Materials Characterization, 1st ed., Manning

Publications Co., 1992.

[9] W.H. Bragg, M. A, F.R. S, The Reflection of X-rays by Crystals, R. Soc. 17 (1913) 43.

doi:10.1098/rspa.1913.0040.

[10] E. Barsoukov, J.R. Macdonald, eds., Impedance Spectroscopy; Theory, Experiments and Applications, 2nd ed.,

John Wiley & Sons, Inc., USA, 2005. doi:10.1002/0471716243.

[11] A.L. Thomann, N. Semmar, R. Dussart, J. Mathias, V. Lang, Diagnostic system for plasma/surface energy

transfer characterization, Rev. Sci. Instrum. 77 (2006) 33501. doi:10.1063/1.2166467.






(2013). doi:10.1016/j.tsf.2013.07.025.

[14] P.A. Cormier, A. Balhamri, A.L. Thomann, R. Dussart, N. Semmar, J. Mathias, R. Snyders, S. Konstantinidis,

Measuring the energy flux at the substrate position during magnetron sputter deposition processes, J. Appl. Phys.

113 (2013) 13305. doi:10.1063/1.4773103.



82

83

5. Influence of oxygen vacancies on

the phase constitution of zirconia

thin films Zirconia (ZrO2) is a polymorphous material which exists in three crystallographic phases

under atmospheric pressure: (i) the monoclinic phase stable up to ~ 1205 °C; (ii) the

tetragonal phase appears from ~ 1205 °C to 2377 °C; and finally (iii) the cubic phase from

2377 °C to 2710 °C (melting temperature) [1]. As cubic zirconia exhibit superior mechanical

and thermal properties over monoclinic zirconia (discussed in chapter 2), the stabilization of

cubic phase of zirconia at room temperature has always been of great importance. For decades

the stabilization has been achieved by the doping the zirconia lattice with cations of lower

valence than Zr (e.g. Y) [2,3] and the resulting material is known as yttria stabilized zirconia

(YSZ) [3]. As doping has its own disadvantages, therefore, to stabilize the high temperature

c-phases of zirconia at room temperature without any doping of yttria, an intense research

has been developed during the last one and a half decade using various synthesis techniques.

The stabilization procedure has been related to the grain size, energy input during growth,

stresses in the film and O vacancies/N atom incorporation in the zirconia lattice [4–14].

However, a consensus over what drives the phase formation in zirconia has not been reached

so far.

84

In this Chapter, we demonstrate that O vacancy incorporation is the sole responsible for the

stabilization of the high temperature c-phase of zirconia, at room temperature. To achieve

this, we coupled cold plasma-based reactive magnetron sputtering experiments to quantum

chemistry based Density Functional Theory (DFT) calculations, performed in collaboration

with Dr. D. Cornil and Prof. J. Cornil, Service de Chimie des Matériaux Nouveaux,

University of Mons.

5.1 Phase stability of oxygen deficient zirconia; quantum chemistry based

DFT calculations

To investigate the influence of oxygen vacancy incorporation on the zirconia phase stability,

O vacancies were introduced in the zirconia lattice randomly, apart from each other and in

the form of clusters (details in chapter 4). Based on DFT calculations, the impact of oxygen

vacancy concentration and their distribution in the zirconia super cells is quantified in terms

of energetics of the ZrO2-x t- and c-phases7. In figure Fig. 5. 1, It is observed that in case of

randomly and apart introduced O vacancies, the c-phase is thermodynamically the most stable

phase if more than 3 at.% of O vacancies are incorporated in the lattice. Similar theoretical

results have also been reported, via a self-consistent tight-binding model, by Fabris et al. [15].

In their publication, these authors suggest that the stabilization of the t- and c-phase of

zirconia can be achieved (only in theory) solely by incorporating O vacancies in the zirconia

lattice. In case of clustered O vacancies, the trend is not clear and is due to the clustering of

O vacancies which result in big holes in the built cell. Finally, we found a very good

agreement from our DFT calculations for the transition enthalpies of ZrO2. The enthalpy for

the m to c transition, equals 14.45 kJ/mol in our case while the one measured by X. Luo et al.

7 Data in numerical form is provided in the Annex I at the end of thesis as well as the energy (eV/atom) is also

presented in the form of graphs for sake of comparison.

85

equals 14.26 kJ/mol [16].

Fig. 5. 1: Influence of O vacancies on zirconia phase constitution. The energies of each phase are

compared as a function of the concentration of oxygen vacancies. Three ways to distribute the

vacancies in the lattice are presented: (a) random, (b) apart from each other, and (c) clustered.

86

5.2 Synthesis of ZrO2-x by reactive magnetron sputtering

To assess experimentally how the incorporation of O vacancies influence the zirconia phase

formation/stabilization 100 ± 10 nm thick ZrO2-x films were deposited inside the transition

zone as well as in the oxidized mode, according to the procedure described in chapter 4. To

deposit the films, Zr target was fed by 200 mA of dc current and the pressure was kept

constant to 10 mTorr. Such deposition parameters were chosen after systematically varying

the deposition parameters.

The chemical composition of the films deposited in such condition is shown in Fig. 5. 2b. It

is observed that the films deposited inside the transition zone at 50, 65, 75, 80 and 100 %

signal (Fig. 5. 2a) are indeed under-stoichiometric and thus contain O vacancies (32, 20, 16,

6, and 3 at.%, respectively). In such under-stoichiometric zirconia, the local Zr charge state

may vary from Zr+ to Zr4+ depending on the surrounding atoms and density of vacancies [17–

21]. Interestingly, the quantum-chemical calculations show that the net charge on zirconium

cation varies from 2.54 |e| for ZrO2 down to 1.87 |e| for 15% of O vacancy. These values are

in good agreement with those reported in [18], which highlight a variation from 2.57 |e| for

ZrO2 down to 2.02 |e| for Zr2O3. According to the DFT calculations presented in Fig. 5. 1 a

and b, such amount of O vacancies should induce the formation of the c-phase. In contrast,

the film deposited in the oxidized mode is stoichiometric. The observed decrease in O

vacancy concentration with the increase in signal (i.e., increase in the target voltage as the

18O2 partial pressure increases) is due to the increased oxide compound formation on the

target surface.

To check the influence of film chemical composition on film crystallinity, grazing incidence

x-ray diffraction (GIXRD) was performed and the resulting diffractograms were compared

87

to ICDD PDF cards representing patterns of the three polymorphs of ZrO2: monoclinic (PDF#

37-1484), tetragonal (PDF# 81-1544) and cubic (PDF# 49-1642). The peak positions of the

films deposited inside the transition zone matched very well with those corresponding to the

tetragonal and/or cubic phases and were hard to distinguish. To unambiguously identify the

phase, we compared the theoretical diffractograms of the t- and c-phases (obtained from the

DFT optimized structures) with the diffractogram of one of our film grown at 100% sensor

signal i.e. film with the lowest concentration of vacancies, as shown in Fig. 5. 3. Furthermore,

Lamas et al [22] have pointed out the splitting of the (400) peak of the c-phase of ZrO2 into

the (004) and (400) peaks in the t-phase of ZrO2, with more than one degree of separation in

their XRD spectra. In Fig. 5. 3, the same peak splitting can be clearly seen for the theoretical

diffractogram of the t-phase while we did not observe any splitting of the (400) peak of the

film grown at 100% sensor signal, thus confirming that the oxygen vacancy-doped films (i.e.

deposited inside the transition zone) belong to the cubic zirconia system.

Fig. 5. 2: (a) Target voltage curve of Zr target as a function of O2 flow shows the transition zone and

working points inside the transition zone as well as in the poisoned zone where (b) chemical analysis

data reveal that ZrO2-x and ZrO2 films were deposited, respectively.

88

The diffractograms of the DFT structures obtained for the various O vacancy concentrations

are shown in Fig. 5. 4(a) together with the diffractograms of the as-deposited films recorded

using an incidence angle of 0.5° (Fig. 5. 4(b)). The remarkable agreement between theory

and experiments demonstrates that the oxygen-vacancy doped films (synthesized inside the

transition zone) are of pure c-phase while the film changes to phase pure monoclinic when

deposited in the oxidized mode, i.e. when containing no O vacancies. However, in DFT based

diffractograms Fig. 5. 4(a), it is observed that some new peaks start to appear with the

increase in O vacancies. The latter could be an artifact coming from the large number of

vacancies in the cell. On the other hand, it is also possible that these peaks are not detected

in the experimental diffractograms because of a too low signal-to-noise ratio.

Fig. 5. 3: Comparison of the diffractograms of the cubic and tetragonal phases obtained at the DFT

level with the experimental diffractogram of the film grown at 100% sensor signal inside the transition

zone (vacancy concentration = 3 % at.). The inset data shows that there is no splitting of the

experimental c(400) peak, as it is observed for the tetragonal (004) and (400) peak.

It has been also reported that incorporating O vacancies induces lattice distortions and result

in contraction of the zirconia lattice, which may result in similar zirconia lattice as in cubic

zirconia due to the coulomb forces between O vacancy-Zr and O vacancy-O atoms [23,24].

89

Furthermore, Fabris et al. [15] have also shown in their work that having a low content of O

vacancies, i.e. 1 at.% (similar to 3.2 mol.% Y2O3), leads to the tetragonal distortion, while

having a higher content of O vacancies, i.e. 4 at.% (=14.4 mol.% Y2O3) results in every

Fig. 5. 4: GIXRD diffractograms of the (a) structures resulting from the DFT calculations

and (b) deposited films, as a function of concentration of O vacancies.

oxygen atom to be a neighbor of a vacant site or at least four of its six neighboring oxygen

atoms, leads to the cubic structure. Moreover, as we do not observe any peak shift in the XRD

90

spectra of the deposited films, is an indication that the deposited films are not stressed. For

these reasons, we believe that the only possible atomistic mechanism behind stabilization of

c-phase of zirconia is lattice distortion caused by O vacancy incorporation in the zirconia

lattice, ultimately forcing the O and Zr atoms to arrange themselves in the high symmetry

cubic crystal.

5.3 Conclusion

In conclusion, based on our DFT calculations and their remarkable agreement with our

experimental data, we conclude that incorporating (as low as 3 at.%) O vacancies is the sole

mechanism responsible of promoting zirconia to the high-temperature c-phase, at room

temperature. In this case, the film deposition was performed with a discharge current of 200

mA, and total pressure of 10 mTorr. However, we observed that deviation from this

“optimized” growth condition (200 mA, 10 mTorr) influence the zirconia phase constitution.

Studying the influence of the working conditions is the subject of chapter 6.

91

References

[1] J.P. Abriata, J. Garcés, R. Versaci, The O-Zr (Oxygen-Zirconium) system, Bull. Alloy Phase Diagrams. 7 (1986)

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[11] F. Lu, J. Zhang, M. Huang, F. Namavar, R.C. Ewing, J. Lian, Phase Transformation of Nanosized ZrO2 upon

Thermal Annealing and Intense Radiation, J. Phys. Chem. C. 115 (2011) 7193–7201.



doi:10.1088/0957-4484/18/41/415702.

[13] Z. Chen, G. Ji, N. Prud’Homme, A. Addad, V. Ji, J. Chevalier, G. Bernard-Granger, Stabilization of the

tetragonal phase in large columnar zirconia crystals without incorporating dopants, Scr. Mater. 68 (2013) 559–

562. doi:10.1016/j.scriptamat.2012.11.028.

[14] M.B. Ponnuchamy, A.S. Gandhi, Lattice expansion and contraction in nanocrystalline yttria-stabilized zirconia

powders, Scr. Mater. 83 (2014) 21–24. doi:10.1016/j.scriptamat.2014.03.028.


Acta Mater. 50 (2002) 5171–5178.

[16] X. Luo, W. Zhou, S. Ushakov, A. Navrotsky, A. Demkov, Monoclinic to tetragonal transformations in hafnia and

zirconia: A combined calorimetric and density functional study, Phys. Rev. B. 80 (2009) 134119.

doi:10.1103/PhysRevB.80.134119.

[17] A. Roustila, A. Rabehi, M. Souici, J. Chene, XPS and AES Study of Oxygen Interaction on the Surface of the

ZrNi Intermetallic Compound, Adv. Mater. Res. 445 (2012) 709–713.

doi:10.4028/www.scientific.net/AMR.445.709.

[18] K.-H. Xue, P. Blaise, L.R.C. Fonseca, Y. Nishi, Prediction of semi-metallic tetragonal Hf2O3 and Zr2O3 from

first-principles, Phys. Rev. Lett. 110 (2013) 65502. doi:10.1103/PhysRevLett.110.065502.

[19] C. Morant, J.M. Sanz, L. Galan, Ar-ion bombardment effects on ZrO2 surfaces, Phys. Rev. B. 45 (1992) 1391–

1398. doi:10.1103/PhysRevB.45.1391.

[20] C. Morant, A. Fernandez, A.R. Gonzalez-Elipe, L. Soriano, A. Stampfl, A.M. Bradshaw, J.M. Sanz, Electronic

structure of stoichiometric and Ar+-bombarded ZrO2 determined by resonant photoemission, Phys. Rev. B. 52

(1995) 11711–11720. doi:10.1103/PhysRevB.52.11711.

[21] B. Puchala, A. Van Der Ven, Thermodynamics of the Zr-O system from first-principles calculations, Phys. Rev.

B - Condens. Matter Mater. Phys. 88 (2013) 94108. doi:10.1103/PhysRevB.88.094108.

[22] D.G. Lamas, R.O. Fuentes, I.O. Fábregas, M.E. Fernández De Rapp, G.E. Lascalea, J.R. Casanova, N.E. Walsöe

De Reca, a. F. Craievich, Synchrotron X-ray diffraction study of the tetragonal-cubic phase boundary of

nanocrystalline ZrO2-CeO2 synthesized by a gel-combustion process, J. Appl. Crystallogr. 38 (2005) 867–873.

doi:10.1107/S0021889805025343.

[23] D.M. Ramo, a L. Shluger, Structure and spectroscopic properties of oxygen divacancy in yttrium-stabilized

zirconia, J. Phys. Conf. Ser. 117 (2008) 12022. doi:10.1088/1742-6596/117/1/012022.

[24] S. Kasamatsu, T. Tada, S. Watanabe, Comparative study of charged and neutral oxygen vacancies in cubic

zirconia from first principles, Appl. Phys. Express. 2 (2009) 2–5. doi:10.1143/APEX.2.061402.

92

93

6. Influence of the deposition

parameters on the phase

constitution of oxygen vacancy

doped thin films

In the previous chapter 5, the stabilization of the high temperature c-phase of zirconia at room

temperature is attributed to the incorporation of O vacancies in the zirconia lattice and is

demonstrated for 200 mA and 10 mTorr sputtering condition. A deviation from the

experimental condition (200 mA and 10 mTorr) for which the diffractogram only exhibits the

cubic reflections, lead to the change in zirconia phase constitution. Moreover, the results

presented in the previous chapter 5 are related to films of thickness ≈ 100 nm. In the

literature, it is shown that the film thickness may also influence the film crystallinity,

microstructure, mechanical and optical properties [1–5]. In similar way one can wonder what

would happen in the case of OVSZ thin films if their thickness is increased. Therefore, to

investigate the influence of such deposition parameters, this chapter is divided into two

sections; i) the evolution of the X-Ray diffractograms is presented as a function of the

deposition pressure and discharge current. The evolution of the XRD data is discussed in

94

terms of energy transferred to the substrate – film system during growth. ii) influence of film

thickness on OVSZ films is investigated by varying the films thickness but the film

deposition conditions were kept constant during the whole deposition process.

6.1 Influence of pressure and discharge current

6.1.1 Experimental details

The deviation from the established sputtering condition is carried out by systematically

varying the deposition condition i.e. by working in the transition zone with the help of the

voltage feedback control unit at 5, 10 and 20 mTorr using 200, 300 and 400 mA of discharge

current, fed to the 5 cm in diameter Zr cathode. In this section, the film characterized by 16

at% O vacancy concentration for all above mentioned working condition and 3 at.% O

vacancy concentration, deposited only at 200 mA, 5, and 10 mTorr are discussed. During the

deposition, the substrates were placed at a distance of 6.5 cm from the target surface.

As the pressure and discharge current can have a direct influence on the energetic of the

plasma species, the global energy flux is measured using the heat flux microsensor developed

by the research group GREMI of the University of Orléans [6–8]. The HFM measures the

contribution emanating from the bombardment of Ar+ and other ions, fast particles such as

O- or secondary electrons emitted from the target, plasma e-, hν, IR emitted by the hot

sputtered target, condensing atoms, etc. During each of these deposition conditions, the heat

flux probe was placed at the substrate position. These measurements were performed at

ChIPS lab, in the frame of a one-week exchange program between ChIPS and GREMI

laboratories. Once the global energy flux was recorded by the heat flux probe, then it was

normalized with respect to the deposition rate i.e. the number of film-forming Zr atoms

reaching the surface of the sample as determined from Rutherford Backscattering

95

Spectroscopy (RBS). More technical details on thin film deposition, GIXRD, energy flux

measurement and on RBS can be found in chapter 4.

6.1.2 Results and discussion

6.1.2.1 Evolution of the phase constitution as a function of pressure and

discharge current

GIXRD performed on the 16 at.% O vacancy films deposited for each of the above mentioned

deposition conditions revealed that working at 10, 20 mTorr using a discharge current 200,

300 and 400 mA lead to the formation of c-phase (Fig. 6.1a). On the other hand, using a

discharge current of 300 and 400 mA and working at low pressure of 5 mTorr, lead to the

mixed c- and m-phase formation. However, working at 200 mA with a pressure of 5 mTorr

resulted in a grey area of phase formation i.e. no clear evidence of pure cubic or mixed c- and

m-phase. All the films were deposited at the same position inside the transition zone i.e. at

the same target coverage. Therefore it is believed the films possess the same film chemistry

i.e. the same O vacancy concentration when they are grown using the same set point value

[9]. It should be noted here that the O vacancy concentration was determined for the case of

200 mA, 5 and 10 mTorr using 18O and NRA/RBS analysis. The vacancy concentration was

found to be identical i.e. 16 at.%, corresponding to a set-point value for the voltage feedback

control unit equal to 75 %. Similarly, NRA/RBS analysis of the films deposited using a set-

point 100% at 200 mA, and pressure 5 and 10 mTorr exhibited the identical concentration of

O vacancy i.e. 3 at.%. GIXRD of 3 at.% O vacancy film revealed film deposited at 200 mA,

10 mTorr is pure cubic, while film deposited at 200 mA, 5 mTorr exhibit a mixed c- and m-

phase (Fig. 6. 1b).

96

Fig. 6. 1: GI-XRD (Cu Kɑ1) diffractograms of zirconia thin films containing (a) 16 at.% O

vacancies, deposited at various discharge currents and pressures. (b) 3 at.% O vacancies.

97

According to the above mentioned observations, one could argue that decreasing the working

pressure and/or increasing the discharge current would increase the energy deposited on the

substrate surface during film growth. Moreover working at higher set point (100% set point

in order to synthesize 3 at.% O vacancy concentration films) inside the transition zone for the

same pressure and discharge current, lead to the higher target voltage as compared to the

lower set-point, 75% signal (16 at.%), see Fig. 5. 2. Increasing the current leads to the

production of more numerous ions and electrons, decreasing the pressure results in an

increase of the mean free path and increase in set-point i.e. voltage will lead to the increase

in energy of the depositing ions. Therefore, to study the influence of deposition parameters

on the global energy flux arriving at the substrate, and ultimately on the film phase

constitution, the energy flux was measured using the heat flux probe. To allow the

comparison between the various deposition conditions, the energy flux was normalized by

dividing the flux value by the flux of deposited Zr atoms, as measured by RBS.

6.1.2.2 Normalized energy flux measurements

The normalized energy flux data are presented in Table 6. 1 for each working condition.

Table 6. 1: Global energy flux per deposited Zr atom (eV/Zr atom) at the substrate position is

reported for various dc-RMS working conditions for the deposition of ZrO2-x thin films containing

16 and 3 at.% O vacancies. It has to be noted the values for 300 mA for 5, 10, 20 and for 400 mA,

10 mTorr are extrapolated values.

16 at.% O vacancy concertation

5 mTorr (eV/Zr) 10 mTorr (eV/Zr) 20 mTorr (eV/Zr)

200 mA 1031 598 395

300 mA 1314 762 504

400 mA 1598 927 613

3 at.% O vacancy concentration

200 mA 1313 823 n.a

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Normalized energy flux data reveal that low energetic values are obtained for a given

discharge current at 20 mTorr working pressure. The normalized energy flux increases

significantly with the decrease in working pressure and reached the keV range on lowest

working pressure i.e. at 5 mTorr. This increase in normalized energy with the decrease in

pressure is primarily related to the increase of the flux and kinetic energy of film-forming

species as the mean free path becomes longer. Other phenomena could also contribute in the

increase of normalized energy flux in this condition, such as the production of fast negative

ions emitted from the oxidized target [10,11]. The later are accelerated in the target sheath

and, in the case of a lower pressure condition, as the target voltage is increased significantly

i.e. 260 V for 5 mTorr and 210 V for 20 mTorr (and thus the kinetic energy of the negative

oxygen ions emitted from the target too) to keep the current constant.

Fig. 6. 2: Influence of energetics on the phase constitution of zirconia thin films containing 16 at.%

O vacancies.

Plotting the resulting film phase as a function of the normalized energy (shown in Fig. 6. 2)

reveal that critical energy level (i.e. ≺ 1000 eV/Zr atom) under which the c-phase is formed

99

could be determined (white zone in Fig. 6. 2). Above 1050 eV/Zr atom (red zone in Fig. 6.

2), the formation of m-phase occurs as well although the film is synthesized in the transition

zone. In between 1000-1050 eV/Zr atom is the grey area with no clarity of phase pure cubic

or mixed c- and m-phase formation. It should be noted that the presence of a critical value of

the normalized energy flux has also been reported by Mraz et al [12] in the case of reactively

sputtered titanium dioxide films. In their study they show that below 540 eV/Ti atom only

the reflections from the anatase phase are obtained. While above 1000 eV/Ti atom, only the

rutile phase appears. In between these two values, a mixture of anatase – rutile phases is

obtained. Cormier et al. [13] have also reported normalized energy flux data in the case of

reactively sputtered titanium dioxide films. In their case it is found that, having ≈7 keV/Ti

atom, the films are composed of anatase phase while for an energy range of 7 – 13 keV/Ti

atom, a mixture of anatase – rutile phase is obtained. Similarly, the influence of energy on

zirconia phases has also been reported by Goedicke et al.[14] in 2000, for films of thickness

equal to 100-150 nm and synthesized by reactive pulsed magnetron sputtering (PMS) at

various deposition pressure. Goedicke et al. also varied the working pressure from 2.2 mTorr

to 26.2 mTorr as well as the target to substrate distance. The sputter power was kept constant.

In their study, when the films were deposited at low pressure i.e. at 2.2 mTorr, the

diffractograms only exhibited the monoclinic lines. On the other hand, when films were

deposited at 26.2mTorr, they exhibited only the cubic peaks. These results are similar to ours,

however, these authors do not specify the elemental composition of their films. They

controlled the oxygen gas inlet by an optical emission monitor but it is not specified if they

work in the transition zone, and they didn’t analyze the XRD data in depth to verify if they

synthesized the tetragonal or the cubic phase. The reason Goedicke et al. provide for such

behavior is related to the energy of the condensing particles. They suggested that the use of

100

higher sputtering pressure reduces the mean free path of the sputtered particles and therefore

the mobility of the depositing species on the substrate. They further relate this lower mobility

to the microporosity of the film and therefore to the reduced hardness and lower compressive

stresses in the film. Goedicke et al. also measured residual stresses and found that, indeed,

films deposited at low sputtering pressure (2.2 mTorr) exhibit high compressive stresses (~ -

1800 MPa) while the films deposited at higher sputtering pressure (26.2mTorr) exhibit low

tensile stresses (~139 MPa). The generation of compressive stresses at low sputtering

pressure could be explained by e.g. the energetic species which get implanted in the growing

film and cause subsurface effects [15]. This implantation leads to lattice defects e.g.,

displacement of atoms from their lattice site. In reactive magnetron sputtering, these high

energetic species could be negatively charged oxygen ions (O-) formed at the oxidized

fraction of the target surface [10,16–18]. As already specified above, these O- ions are

accelerated in the cathode sheath and bombard the growing film with energies in the range

of several hundreds of eV [10]. The energy of these O- ions depend on the magnitude of the

target voltage, while the number of these O- ions is determined by the target coverage by the

oxide layer. In several studies, the influence of these O- ions on the structure formation has

been reported. For example, in 2006, Mraz et al.[11] studied the influence of the these oxygen

ions emitted from the oxidized target during the growth of Nb, Ta, Zr, and Hf oxide films by

reactive magnetron. These authors proposed that the evolution of the crystalline structure of

the transition metal oxide may depend on the presence of O− ion bombardment induced

adatom mobility. Ngaruiya et al. [18] also studied the structure formation of various transition

metal oxides of group 4, 5 and 6 by reactive magnetron sputtering at lower pressure (6

mTorr). Ngaruiya et al. found that Zr and Hf-based sputtering processes allow the formation

of the monoclinic phases of their respective oxides. To the contrary, the other transition metal

101

oxides from group 5 and 6 (Nb, Ta, W, V, and Mo) form amorphous films in the same

deposition conditions. For the zirconium target, Ngaruiya et al. observed that monoclinic

zirconia films were deposited with high deposition stresses (-1500 MPa) in the case of a fully

oxidized target. They attributed the generated compressive stresses in the film to the flux of

oxygen ions emitted from the oxidized target. In the work of Severin [19], nitrogen is added

to the Ar/O2 atmosphere to reduce the flux of fast negative O- emitted from the target. If the

flow of nitrogen is not high enough, the film phase constitution is monoclinic. On the other

hand, when enough nitrogen is added, the phase is cubic. Actually, by adding N2, the reactive

discharge can be stabilized in the transition zone. Later on, Sarakinos et al [20] discussed the

phase formation of hafnium oxide (which is isostructural with ZrO2) by reactive sputtering.

They confirmed that the suppression of the O- bombardment and the incorporation of oxygen

vacancies favors the generation of cubic crystals but they incorporated the vacancies by using

a High-Power Impulse Magnetron Sputtering (HiPIMS) discharge which is known to increase

the ion bombardment of the film during growth [21].

From the above mentioned results, it appears that the presence of energetic species (such as

fast O- ions) increases the value of the energy deposited during growth and the stress levels

in the coatings. This variation in the growth conditions may induce some modifications in

the film phase constitution. From the data displayed in Fig. 6. 1, and 2, it could be speculated

that in our condition, and despite the incorporation of vacancies, the appearance of the

monoclinic phase is the result of the increased energy flux above a certain threshold, which

leads to the build-up of large compressive stress. However, to unambiguously pinpoint the

effect of such energetic species a more detailed investigation must be carried out. Moreover,

as the presented normalized energy flux is a combination of various plasma species

102

contribution therefore, it would be interesting to develop a strategy to figure out the role of

each specie in the phase formation.

6.2 Influence of film thickness on the crystal structure of OVSZ films

6.2.1 Experimental details

OVSZ films with 16 at.% O vacancies were deposited by working inside the transition zone

at 10 mTorr and 200 mA (i.e. the optimized conditions) of discharge current using dc reactive

magnetron sputtering. In order to investigate the influence of film thickness on film

crystallinity, film thickness was varied from 150 nm to 1300 nm. To investigate the influence

of the film thickness on the film crystallinity, GIXRD as well as Bragg-Brentano XRD was

performed on the deposited samples. High resolution transmission electron microscopy (HR-

TEM) was also performed on the deposited samples and for HR-TEM analysis, 2000 nm

thick sample with 16 at.% O vacancies was used and the sample cross section preparation

(sandwich gluing, mechanical polishing and ion polishing) was carried out at -60 °C to avoid

any sample degradation. TEM analysis were carried out at the Institut des Materiaux de

Nantes (Dr. A.A. El Mel and E. Gautron). To investigate the film microstructure, scanning

electron microscopy (SEM) was performed on the cross sections of the 2000 nm thick films.

More technical details regarding the film deposition, HR-TEM, XRD can be found in chapter

4.

6.2.2 Results and discussion

6.2.2.1 Evolution of XRD diffractograms as a function of film thickness

GIXRD and Bragg-Brentano XRD diffractograms of OVSZ films whose thickness equals

150, 275, 615 and 1300 nm are shown in Fig. 6. 3 (a) and (c), respectively. It has to be noted,

as the thicker (615 and 1300 nm) films have more material volume as compared to thinner

103

(150 and 275 nm) films, giving higher intensity of the XRD peaks as compared to thinner

ones. Therefore, to avoid any peak suppression, a narrow scan in the range of 26 - 33° was

also performed on 150 and 275 nm thin films with a 10 time more acquisition time as

compared to 1300 nm thick film and no additional peaks were observed in this 26 - 33° range

Fig. 6. 3: (a) GIXRD, (b) GIXRD narrow scan with 10 time more acquisition time, (c) Bragg-

Brentano diffractograms of 16 at.% O vacancy containing zirconia films of various thickness.

104

(diffractograms shown in Fig. 6. 3b). The first feature that can be noticed with the increase

in film thickness from the GIXRD as well as Bragg-Brentano spectra (shown in Fig. 6. 3) is

the competition between the c(111) and the c(200) peak intensities. It is observed that the

c(111) orientation dominates at lower film thickness i.e. at 150 nm, while with the increase

in film thickness i.e. around 600 nm, the c(200) orientation overtakes the c(111) peak

intensity. Similar feature has also been observed by Nouveau et al. [5] in Bragg-Brentano

XRD diffractograms for CrN films which have a similar cubic crystal structure as cubic

zirconia. Nouveau et al. showed that the compressive stress varies with respect to the film thickness

[5,22] and they observed three growth regimes : Regime A) (up to 150nm) where the compressive

stress keeps increasing, B) (150-500nm) the film changes its structure i.e. preferred plane orientation

to release compressive stress, C) (above 500nm) the compressive stress value is stable and relaxation

is induced by the growth of less dense planes. Another feature can be observed in Fig. 6. 3: the

film with the lowest thickness (150 nm) exhibit pure c-phase, however with the increase in

film thickness a small bump appears around 28°. This is a reflection related to the monoclinic

phase of zirconia. At 1300 nm, this peak associated with the appearance of the monoclinic

phase is significant.

6.2.2.2 Analysis of film cross-section by SEM and HRTEM

To investigate the appearance of monoclinic peak, SEM cross-sectional images and HR-TEM

analysis were performed on 2000 nm thick OVSZ film. Fig. 6. 4 shows the cross-section

images taken at (a) middle, (b) near the top, (c) middle of film cross-section and (d) film

substrate interface reveal that, near the interface the microstructure of the film indicates a

dense nucleation zone (Fig. 6. 4d) followed by columnar feather-like growth (Fig. 6. 4 (b)

and (c)).

105

Fig. 6. 4: SEM cross-sectional images of 2000 nm thick zirconia film having 16 at.% O vacancies (a)

middle, (b) near the top, (c) middle of film cross-section and (d) film-substrate interface. For all these

cross-sectional images small white bar at the right bottom of images correspond to the scale bar which

is 100 nm.

Fig. 6. 5: Cross sectional image of the film and the substrate as obtained by TEM. The substrate is

Si (100). Colored circles indicate the area where selected area electron diffraction analysis was

carried out.

106

Fig. 6. 5 shows the cross-section of the 16 at.% O vacancy film of thickness 2000 nm. To

perform the SAED (Selected Area Electron Diffraction) analysis along the substrate – to –

film surface axis, different areas were selected from the film-substrate interface, as shown in

Fig. 6. 5. SAED analysis, close to the substrate, revealed that the film is textured and belongs

to cubic structure. This area of film could be related to the 150 nm and 275 nm-thick films

from which the XRD analysis which revealed the same cubic structure (Fig. 6. 3). SAED

analysis performed while moving away from the film-substrate towards the middle of the

film also revealed the presence of cubic crystals. No monoclinic crystals were detected as

observed in XRD diffractograms (Fig. 6. 3). The film away from the interface has feather-

like columnar structure as shown in Fig. 6. 4. To locate the monoclinic crystals, a much more

careful investigation was carried out at the intersection of those feathers, as shown by the

yellow rectangle in Fig. 6. 6.

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Fig. 6. 6: HR-TEM images of 16 at.% OVSZ film at around 850 nm from substrate.

At the intersection of two feathers, due to the superposition of feathers, it was hard to

distinguish if the phase is amorphous or not. To avoid this, Fast Fourier Transformation (FFT)

was used. FFT analysis revealed there exists crystals which could only be indexed by

monoclinic phase (Fig. 6. 7). Furthermore, the monoclinic crystals were localized around 850

nm and 2000 nm away from the film-substrate interface hence confirming the initial XRD

analysis (Fig. 6. 3). The above mentioned results highlight that, despite the film growth

conditions are kept constant during the whole deposition process, the film microstructure and

phase constitution may evolve.

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Fig. 6. 7: Electron diffraction of 16 at.% OVSZ film obtained at around 850 nm from the substrate-

film interface. The diffraction spots are indexed with monoclinic phase.

6.3 Conclusions

In the section 6.1 we have presented that oxygen vacancy doped cubic zirconia coatings are

actually obtained in some specific working conditions i.e. low current, moderate to high

pressure. As the current is increased and/or the pressure decreased, monoclinic peaks appear

on the diffractograms, although the films are deposited inside the transition zone. The

evolution of the film phase constitution with the deposition parameters is interpreted in terms

of energy flux at the substrate surface. Pure cubic oxygen vacancy doped zirconia thin films

are obtained by working in conditions for which the normalized energy flux lies below ~1000

eV/Zr atom. Above this threshold, up-to 1050 eV/Zr atom the diffractograms exhibit the grey

area with no clarity of phase pure cubic or mixed c- and m-phase formation but, above 1050

109

eV/Zr atom both cubic and monoclinic peaks. The increase in film stress as a result of the

increased normalized energy flux is proposed to explain this behavior.

In the section 6.2, 16 at. % OVSZ film with various thicknesses, ranging from 150 to 2000

nm, have been deposited. GIXRD diffractograms reveal the appearance of monoclinic peaks

with the increase in film thickness, while the films are deposited in the transition zone. In the

case of thicker films, SEM cross-sectional images show dense nucleation near the substrate-

film interface followed by columnar feather-like growth of the film away from the substrate-

film interface. TEM analysis of these thick films allows to determine that the monoclinic

crystals indeed appear in the upper part of the film, e.g. above 850 nm. The reason for the

appearance of the monoclinic crystals close to the film surface is not yet clarified.

Nevertheless, an important information can be extracted from these results. One should pay

attention when synthesizing (thick) films because, even though the deposition conditions are

kept constant throughout the deposition process, the film microstructure and its phase

constitution may vary.

110

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antimony doped SnO2 thin films, Thin Solid Films. 591 (2015) 18–24. doi:10.1016/j.tsf.2015.08.013.

[2] R. Mariappan, V. Ponnuswamy, P. Suresh, N. Ashok, P. Jayamurugan, A. Chandra Bose, Influence of film

thickness on the properties of sprayed ZnO thin films for gas sensor applications, Superlattices Microstruct. 71

(2014) 238–249. doi:10.1016/j.materresbull.2010.06.041.

[3] D.A. Hardwick, The mechanical properties of thin films: A review, Thin Solid Films. 154 (1987) 109–124.

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[4] A.J. Detor, A.M. Hodge, E. Chason, Y. Wang, H. Xu, M. Conyers, A. Nikroo, A. Hamza, Stress and

microstructure evolution in thick sputtered films, Acta Mater. 57 (2009) 2055–2065.

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[5] C. Nouveau, M. Djouadi, O. Banakh, R. Sanjinés, F. Lévy, Stress and structure profiles for chromium nitride

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7. Thermal stability of OVSZ thin

films

Zirconia is one of the materials which exhibit very low thermal conductivity (1.2 - 2.6

Wm−1K−1[1]) making it a good candidate for thermal barrier coatings (TBCs)[2]. However,

the use of zirconia in high temperature applications is restricted by the change in its volume

(~5 Vol.%)[3] due to phase transformation at elevated temperatures as well as while cooling

down the device. This phase transformation and change in volume result in the crumbling of

the zirconia-based components. Since last couple of decades, phase transformation of

zirconia upon thermal cycling is avoided by stabilizing the high temperature c-phase of

zirconia. Typically, stabilization is achieved by doping zirconia with cations of lower valence

than Zr e.g., Y (as discussed in chapter 2). In such stabilization process some Zr4+ cations are

replaced by Y3+. in this situation, to maintain the charge neutrality, one oxygen vacancy is

created for each two substituting Y3+ cations. This makes YSZ not only useful for TBCs but

the presence of vacancies also makes it a good ionic conductor used as an electrolyte in solid

oxide fuel cells (SOFC)[4–6] and in oxygen sensors[7].

In chapter 5, it is demonstrated that the zirconia high temperature c-phase can be stabilized

at room temperature by incorporating the right amount of oxygen vacancies in the lattice [9].

This material was then called oxygen vacancy stabilized zirconia (OVSZ). By definition,

112

OVSZ contains vacancies and, like YSZ, it could be useful as ion conductor in Solid Oxide

Fuel Cell.

In this chapter, the thermal stability and structural transformation at elevated temperature is

investigated using temperature-resolved XRD measurements in order to verify if, indeed,

OVSZ thin films can be used in high-temperature applications.

7.1 Experimental details

To investigate the thermal stability of the OVSZ coatings, two identical sets of 100±10 nm

thick c-ZrO2-x films containing ~16 and 3 at.% O vacancies were deposited according to the

procedure described in chapter 4. The samples were deposited on Si(100) substrates whose

thickness is 525 ±20 μm (Siegert Wafer).

To analyze the film crystallinity of the as-deposited films, grazing incidence x-ray diffraction

(GIXRD) is performed using PANalytical Empyrean with a Cu-Kα1 radiation source (λ =

0.15406 nm). The diffractograms were recorded with a step size of 0.07° with a duration of

15 seconds using an incidence angle of 0.5° at 40 mA, 45 kV of generator settings.

Temperature-resolved XRD measurements are performed on a Bruker D8 Discover

diffractometer, using a Co-K⍺ radiation (λ = 0.17902 nm) and a parallel beam with a diameter

of 1mm. The measurements were carried out at the Institut Jean Lamour, in Nancy (Pr. J.F.

Pierson and Dr. P Boulet). Samples were heated either in air or in a N2 ambient with a Domed

Hot Stage (DHS1100 from Anton Paar), installed on the goniometer. The heating speed was

8 °C/s. During annealing in the N2 ambient, a flow of 1-2 L/h of N2 was blown at the sample

surface and the pressure inside the dome was same as outside the dome, i.e. 1 atm. These

measurements were carried out in grazing incidence geometry as well but at a grazing angle

of 3° with a step size of 0.025° and duration of 4 s. During the scan, for each working

113

temperature, the annealing stage temperature was kept constant for 2 hours to allow the scan

to complete at that specific temperature. Further, the detection was assured by a scintillator

counter before which long Sollers slits are installed to select the diffracted X-rays.

7.2 Results and discussion

XRD diffractograms of the as deposited films containing 16 and 3 at.% O vacancies acquired

at an incidence angle of 0.5° are shown in Fig. 7. 1. The diffraction peaks from the cubic

phase of zirconia are identified.

Fig. 7. 1: GIXRD spectra of the as-deposited films recorded using Cu x-ray source at an incidence

angle of 0.5° at room temperature.

TR-XRD diffractograms acquired at an incidence angle of 3° for the same films containing

16 and 3 at.% O vacancies annealed in air and N2 from 200 °C to 900 °C and then cooled to

100 °C are shown in Fig. 7. 2. It is observed, for low temperatures (i.e. from 200-300 °C),

that the diffractograms do not show any peaks. Let’s recall that, for these specific

experiments, the films are scanned at an incidence angle of 3° which is a higher value than

for the diffractograms presented in Fig. 7. 1, where the same films where scanned at an

incidence angle of 0.5°. It is known that in GIXRD, the intensity of the peak is very much

dependent on the incidence angle i.e. at low incidence angles, close to the critical angle of

114

115

Fig. 7. 2: TR-GIXRD spectra of the films containing approximately 16 at.% O vacancies annealed in

(a) air and (b) N2. Films containing 3 at.% O vacancies are also annealed in (c) air and (d) N2. X-ray

source Co and x-ray incidence angle 3° was used.

116

total reflection, the beam path inside the probed material is increased (beam path ≈ t/Sinθ, t

is film thickness and θ is beam incidence angle[10]). As a result, an increase in the intensity

of the reflected peak is observed when films are scanned at lower incidence angles. Further

annealing of films at 350 °C shows that the cubic diffraction peaks become more intense,

meaning the films get better crystallized. This could be explained in terms of diffusion of

atoms inside the film material at higher temperature. This phenomenon is supported by the

observed increase in the intensity of c(111) peak as a function of temperature, shown in Fig.

7. 3. This situation allows already existing small cubic crystals to grow in size by gathering

diffusing atoms. At 350 °C those crystals are big enough (average grain size 12-15 nm at 350

°C while average grain size of as deposited films is 8-9 nm) to be detected in the used TR-

XRD setup. The evolution of the crystallite size as a function of the temperature is shown in

Fig. 7. 4. The crystallite size is calculated using the Scherrer Equation [11] and taking the

c(111) peak into consideration. Data show for 16 and 3 at.% Vo films that the crystallite size

increases monotonously with the increase in temperature, up to 700 °C. After passing 700 °C

a relatively fast increase is observed. Further, on cooling down the films after reaching 900

°C, it is observed that, in both cases, the average crystallite size remains around 20 nm.

Fig. 7. 3: Evolution of c(111) and m(111) as function of temperature of film containing 16% O

vacancies.

117

Fig. 7. 4: Evolution of cubic crystallite size as a function of temperature, calculated using C(111)

peak of films containing (a) 16 at. % of O vacancies and (b) 3 at. % of O vacancies.

The analysis of the annealed samples shows that, after passing 750 °C, m(111) peak appears

and gets more and more prominent with the increase in temperature. It is also observed that

nor the annealing ambient (i.e., air or N2) nor the Vo concentration has any influence on the

appearance of the m-peak. However, with the increase in annealing temperature a significant

shift to lower angles in all peak positions is observed on the TR-XRD spectra (however this

peak shift is different for each peak, not shown here), whatever the oxygen vacancy

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Fig. 7. 5: Shift of C(111) peak as a function of temperature of the films containing (a) 16 at.% of O

vacancies (b) 3 at.% of O vacancies to lower angles during annealing in Air and N2 and recovery

while cooling down.

concentration (16 or 3 at.%) or the annealing ambient is (air or N2). At 350 °C, the c(111)

peak exhibits a shift of less than 0.2°. With the increase in annealing temperature, the c(111)

peak shifts further to lower angles due to the thermal expansion. At 900 °C, a maximum shift

to lower angle of about 0.5° is reached as shown in Fig. 7. 5. While on cooling down the films

it is observed that the c(111) peak tries to return to its original position. One should also note

that the m(111) peak also did not vanish upon cooling the films to room temperature (shown

119

in Fig. 7. 2). It is also noted from Fig. 7. 2 and Fig. 7. 5, that the films containing 16 and 3

at.% Vo and annealed in air and N2 show a similar evolution.

The more detailed analysis of the temperature–resolved diffraction data of 16 and 3 at.% Vo

reveals that, after passing 750 °C, the area under the c(111) start to decrease (Fig. 7. 6). On

the other hand, at this precise temperature, the m(111) peak appears as shown in Fig. 7. 2.

The variation of the area under the m(111) peak as a function of temperature is almost

identical (in absolute value) to the decrease in area under the c(111) peak, as shown in Fig.

7. 6 for 16 at.% and 3 at.% Vo. Such a decrease in area under the c(111) peak after passing

750 °C would highlight the destruction of cubic zirconia crystals. To the contrary, monoclinic

crystals are formed. From the respective variation in intensities, one could, at first sight,

believe that all the monoclinic crystals are coming from the transformation of the cubic

crystals after 750 °C. Here it should be highlighted that the diffraction intensities related to

the c(111) and m(111) reflections are not identical when looking at the synthetic XRD

spectrum constructed from monoclinic and cubic unit cells, built from quantum chemistry–

based calculations. These calculations and the ZrO2 unit cell utilized to generate the synthetic

XRD spectra are presented in chapter 5. From the synthetic spectra it is found that the

relationship between the diffracted intensities, for a single diffracting unit cell, is Im(111) = 0.35

Ic(111). So having almost the same decrease and gain in area by the cubic and monoclinic peaks

respectively would point out there are more monoclinic crystals formed than cubic crystals

destroyed.

120

Fig. 7. 6: Area under the c(111) and m(111) peak as a function of temperature of film containing 16

at.% O vacancies. After passing the 750 °C a drop in c(111) and increase in m(111) peak area is

observed. Similar evolution is observed for film containing 3 at.% O vacancies.

According to these observations, the appearance of m-phase could eventually be explained

by two mechanisms; i) the oxygen uptake by a small fraction of the annealed OVSZ thin film

from the surroundings after passing 750 °C. That part of the film would therefore turn into

stoichiometric m-ZrO2 and/or ii) the martensitic transformation of existing cubic crystals, as

in case of YSZ. It should be also noted that monoclinic structure is the thermodynamically

stable phase below ~1200°C, for fully stoichiometric ZrO2.

Since the OVSZ thin films are, by definition, under-stoichiometric, the possibility of oxygen

uptake and the annihilation of the oxygen vacancies by the incorporation of O from the

surrounding ambient cannot be ignored. To this this purpose, RBS and NRA measurements

were carried out on the OVSZ thin films (deposited using 18O and a graphite sheet as

substrate) at room temperature and after annealing at 750 °C for two hours. It was observed,

due to the annealing at 750 °C, 18O was completely replaced by 16O. Here, it should be

mentioned that the top view of the films grown on the graphite substrate shows the presence

of cracks already before the annealing process, Fig. 7. 7. The

121

Fig. 7. 7: SEM top view images of as deposited films showing cracks. Films are deposited on

graphite substrate for RBS and NRA analysis.

presence of cracks most likely change dramatically the way the film interacts with the

surrounding atmosphere, and the amount of oxygen it takes up upon annealing. However,

because of some instrumental reasons it was hard to calculate the 16O uptake concentration

from the surrounding ambient. Furthermore, as it has been demonstrated in the previous

chapter 5, that oxygen vacancy incorporation promotes the formation of the c-phase at low

temperature and, as the films still retain some of the c-phase up to 900 °C and even on cooling

down to 100 °C (Fig. 7. 2), it seems not appropriate to claim that oxygen vacancies are

completely annihilated and the films fully turns to stoichiometric ZrO2 at such higher

temperatures (when the film are grown on silicon single crystals). However, it could be that

only a small volume fraction of the film turns to ZrO2. The monoclinic crystals could nucleate

at these particular places. The formation of a stoichiometric monoclinic (passivation) layer

on the film surface would impede the further oxidation of the under-stoichiometric zirconia

left below. In this condition, the diffraction peaks related to the presence of the cubic phase

still remain intense at 900°C and as well as when cooled down to 100 °C, as observed in Fig.

7. 2 and Fig. 7. 6. On the other hand, here one should also note that the evolution of the film

phase constitution as a function of temperature is identical whatever the annealing

122

experiments are performed in air or in the nitrogen ambient. It could be that the latter contains

enough oxygen or water vapor to allow the partial oxidation of the ZrO2-x film.

The 2nd possible mechanism is related to the martensitic transformation of existing cubic

crystals into monoclinic crystals. It is similar to c/t to m phase transformation in sol-gel based

ZrO2 films [12] or stress/pressure induced t to m phase transformation in YSZ [13–15].

Mehner et al. [12] deposited c/t, 1000 – 1200 nm thick, zirconia films on steel substrates

using sol-gel method and studied their thermal stability. Mehner et al. zirconia films were

stable up-to 600 °C and after that temperature, monoclinic peaks appeared on their Bragg-

Brentano diffractograms. They also observed a small shift in the diffraction peak position to

lower angles with the increase in temperature, similar to the shift monitored in the present

study (shown in Fig. 7. 5). Mehner et al [12]attributed the shift to the generation of

compressive stresses in film during annealing. As the buildup of compressive stresses parallel

to the surface causes a vertical expansion of the deposited film and thus results in the increase

of the distance between the planes which are parallel to the substrate surface, the observed

shift in peak to lower angles is an indication of the presence of in-plane compressive stresses

in the films upon annealing. To calculate such annealing generated stresses, Mehner et al.

performed Sin2ψ method on his samples using the c(333) peak. They found below 625 °C

films did not possess high stresses but above 625 °C films exhibit ≈ −750 MPa compressive

stress. They correlated it to the onset of monoclinic phase formation. In the present study,

due to the film thickness (≈ 100 nm), the grazing incidence XRD geometry was used, and in

such geometry the planes probed are not all parallel to the sample surface. Therefore, it’s

hard from the present measurements to calculate the residual stresses from such peak shifts

(Fig. 7. 5). However, since stresses are also generated due to the difference in film-substrate

thermal expansion coefficients (TEC), therefore, a theoretical calculation can be made to

123

estimate the thermal stresses. As Si has a lower TEC (3.6×10-6/K[16]) as compared to c-

zirconia (8.8-10.6×10-6/K[1]) and as the deposited film is bound to the substrate, at elevated

temperature, the film will experience compressive stresses. Such thermal stresses arising

from the mismatch between the film and substrate TEC can be calculated using Eq. 6. 1 where

𝐸𝑓 is the Young’s modulus, 𝑣𝑓 is the Poisson ratio, 𝛼𝑓 and 𝛼𝑠 are the respective TECs of the

film and substrate, 𝑇𝑠 is the substrate temperature before starting the annealing (25°C), and

𝑇𝑎 is the temperature during the measurement. A negative value of 𝜎𝑡ℎ corresponds to

compressive stress.

𝜎𝑡ℎ =𝐸𝑓

1 − 𝑣𝑓(𝛼𝑓 − 𝛼𝑠)(𝑇𝑠 − 𝑇𝑎) (6.1)

To assess the theoretical thermal stress values in such deposited films, the Young’s modulus

and Poisson ratio was taken from the literature (E =200 GPa [17,18], , 𝑣𝑓 = 0.25 [12]) and

we assume that the zirconia Young’s modulus variation with the temperature is negligible.

The calculated compressive stress was found to increase with the increase in temperature as

shown in Fig. 7. 8. Calculation indicates that, at 750 °C, the film stress reaches a value of -

856 MPa. At 900 °C, the value reaches -1032 MPa. Comparing these results with Mehner’s

et al. report (c/t to m phase transformation occurs at ≈ −750 MPa, at) it appears that OVSZ

film are more resilient, i.e. stable up to ≈ −856 MPa, than the sol-gel based zirconia films of

these authors. Moreover, it can be assumed that in such circumstances, where the

compressive stresses are building up inside the film, another probable reason for the

appearance of the m-peak (beside oxygen uptake from the annealing atmosphere) is that some

cubic crystals turn into monoclinic crystals in order to cope with the mechanical stress, i.e.

through a martensitic, diffusion-less, transformation. So, per these arguments, it seems

reasonable to assume that our OVSZ films reached a critical threshold of stress after passing

124

750 °C, and such stresses might force the high symmetric cubic structure to transform in to

the low symmetric monoclinic structure. From our data, this critical threshold value of stress

appears to be in the range of (≈ −856 MPa), as shown in Fig. 7. 8.

As a final word, it should be added that when thermal cycling was performed (6 times in the

present study) below 700 °C, the diffractograms of the OVSZ films only exhibit the c-phase

reflections (data not shown).

Fig. 7. 8: Theoretical residual stress evolution as a function of temperature. Vertical dashed lines

show the temperature where monoclinic peak appeared.

7.3 Conclusion

From the above presented data, we conclude that oxygen vacancy stabilized cubic zirconia

(OVSZ) films deposited on Si wafers are stable up-to 750 °C. It is suggested that a

combination of two mechanisms is responsible for the appearance of m-phase above the

critical temperature of 750 °C; i) oxygen uptake by a fraction of OVSZ thin film from the

surroundings which leads to the formation of stoichiometric ZrO2. The later crystallizes under

the monoclinic structure, thermodynamically stable in such circumstances. ii) The martensitic

transformation of existing cubic crystals, due to the compressive stress forcing the cubic

125

crystals to rearrange themselves in to the low symmetric m-phase. However, to

unambiguously separate these two mechanism a more detailed study is needed.

The development of such high temperature stable c-zirconia film containing oxygen

vacancies might not only be helpful in thermal barrier coating applications but also in SOFC

applications, as ionic conductive electrolyte membranes. This thin film material would allow

to lower the operating temperature of such devices while offering high ionic conductivity.

The analysis of the ion conductivity of the as-deposited OVSZ films is the subject of the next

chapter of this thesis work.

126

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doi:10.1361/10599630418176.

[2] D.R. Clarke, S.R. Phillpot, Thermal barrier coating materials, Mater. Today. 8 (2005) 22–29. doi:10.1016/S1369-

7021(05)70934-2.


Mater. Rev. 50 (2005) 20.

[4] B.C.H. Steele, Oxygen transport and exchange in oxide ceramics, J. Power Sources. 49 (1994) 1–14.

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[9] M. Raza, D. Cornil, J. Cornil, S. Lucas, R. Snyders, S. Konstantinidis, Oxygen Vacancy Stabilized Zirconia

(OVSZ); A Joint Experimental and Theoretical Study, Scr. Mater. 124 (2016) 26–29.

doi:10.1016/j.scriptamat.2016.06.025.

[10] R. Feder, B.. Berry, Seeman-Bohlin X-ray Diffractometer for Thin Films, J. Appl. Crystallogr. 3 (1970) 372–379.

[11] A.L. Patterson, The scherrer formula for X-ray particle size determination, Phys. Rev. 56 (1939) 978–982.

doi:10.1103/PhysRev.56.978.



[13] J.R. Piascik, Q. Zhang, C.A. Bower, J.Y. Thompson, B.R. Stoner, Evidence of stress-induced tetragonal-to-

monoclinic phase transformation during sputter deposition of yttria-stabilized zirconia, J. Mater. Res. 22 (2007)

1105–1111. doi:Doi 10.1557/Jmr.2007.0128.

[14] M. Allahkarami, J.C. Hanan, Residual Stress and Phase Transformation in Zirconia Restoration Ceramics, Adv.

Bioceram. Porous Ceram. V Ceram. Eng. Sci. Proc. 574 (2012) 37. doi:10.1002/9781118217504.ch6.

[15] M. Allahkarami, J.C. Hanan, Mapping the tetragonal to monoclinic phase transformation in zirconia core dental

crowns, Dent. Mater. 27 (2011) 1279–1284. doi:10.1016/j.dental.2011.09.004.

[16] W.M. Yim, R.J. Paff, Thermal expansion of AlN, sapphire, and silicon, J. Appl. Phys. 45 (1974) 1456–1457.

doi:10.1063/1.1663432.

[17] R. Morrell, Handbook of Properties of Technical and Engineering Ceramics, Her Majesty’s Stationary Office,

London, 1985.



127

8. Ionic conductivity of OVSZ thin

films

Use of stabilized zirconia as an electrolyte in SOFCs has always been a top choice due to its

stability at operating temperatures, good ionic conductivity and because of its desirable

chemical stability in both oxidizing and reducing atmospheres [1]. The ionic conductivity in

stabilized zirconia is attributed to the mobility of oxygen vacancies, which are created by the

addition of aliovalent dopants, as discussed in chapter 2. It was first expected that the ionic

conductivity will increase with the increase in O vacancy concentration i.e. by increasing the

dopant content. However, later it was observed the maximum ionic conductivity in yttria-

stabilized zirconia (YSZ) occur at 7 – 9 mol% Y2O3 at 327 – 1227 °C [2,3]. On the other

hand, the incorporation of a higher amount of Y2O3 was found to lower the mobility of O

vacancy by increasing the diffusion energy across an Y–Y common edge as compared to the

diffusion across one with a Zr–Y common edge [2]. In the present chapter, we will measure

the ionic conductivity of Oxygen Vacancy Stabilized Zirconia (OVSZ) thin films in order to

establish if such material can also be utilized as an electrolyte membrane e.g. in SOFC

applications. In the chapter 7, the stability of the magnetron sputtered OVSZ films is already

demonstrated and OVSZ can withstand annealing procedure up to 750 °C.

128

In this section, we measure the ionic conductivity and investigate the influence of O vacancy

concentration, film thickness, and the nature of the substrate on the ionic conductivity of

OVSZ thin films.


In order to perform this study, c-phase OVSZ thin films containing 3 or 16 at.% of O

vacancies and characterized by a thickness of 10, 20, 50 or 100 nm were prepared using dc-

reactive magnetron sputtering, as described in chapter 4. The films were deposited on

polished NdGaO3 (100) substrates (size 5 x 5 x0.5 mm3) purchased from Crystal Germany.

To study the influence of substrate, films containing 16 at.% O vacancies of thickness 10, 20

and 50 nm were also deposited using the same deposition conditions on polished MgO

substrates (size 5 x 5 x0.5 mm3) purchased again from Crystal Germany. To determine the

ionic conductivity, electrochemical Impedance Spectroscopy measurements were carried out,

at the Department of Energy Conversion and Storage, Technical University of Denmark with

the help of Dr. S. Sanna and Dr. V. Esposito), to measure the ionic conductivity at various

temperatures. Four-probe method was also used on the samples containing 3 and 16 at.% O

vacancies of 100 nm thickness deposited on Si (100) and borosilicate glass in order to

measure the film resistivity.


The Arrhenius plot, i.e. the plot of the logarithm of the ionic conductivity σ as a function of

the temperature T, for 10 nm thick OVSZ thin film containing 3 at.% O vacancies is shown

in Fig. 8. 1. According to this figure, the oxygen vacancy doped zirconia thin film is

exhibiting a measurable ionic conductivity in a temperature range which is comparable to

SOFC systems.

129

For the sake of comparison, the logarithm of the ionic conductivity of an YSZ film (thickness

= 15 nm) containing a similar amount of O vacancies (1 at.% O vacancy = 3.2 mol% Y2O3)

and reported by Kosacki et al. [1] is also presented as a function of the temperature on Fig.

8. 1.

Fig. 8. 1: The Arrhenius plot of as deposited OVSZ thin film containing 3 at.% Vo (blue data set) is

compared to an YSZ thin film (red data set, from Kosaki et al [1]) presenting similar physical

characteristics.

The logarithm of the ionic conductivity is proportional to the temperature [4,5]. In both cases,

an increase in ionic conductivity with the increase in temperature is observed (Fig. 8.1). Data

shows that the 3 at.% OVSZ thin film exhibits ~1 S.cm-1 ionic conductivity at 650 °C. The

conductivity of the film rises to ~7.4 S.cm-1 at 725 °C. On the other hand, the YSZ film

exhibits ~ 0.3 S.cm-1 at 725 °C [1]. This is almost 24-fold increase in the ionic conductivity

exhibited by OVSZ as compared to YSZ films.

130

The ionic conductivity 𝜎 can be expressed as 𝜎 = 𝑛𝑞𝜇 , where n is charge carrier

concentration (cm-3), q is the electric charge of the moving ion (here in our case O2- ions) and

μ is the mobility of charge carriers (meaning mobility of O ions) (cm2 s-1 V-1) [6]. As the

charge carrier and their number i.e. O vacancy concentration, in 3 at.% OVSZ thin film and

in a 10 mol% YSZ is almost the same, therefore, this lead to the conclusion that the mobility

of O ions is higher in OVSZ film as compared to YSZ coating because of the high ionic

conductivity measured (~7.4 S.cm-1) at 725 °C. This is also consistent with Pornprasertsuk et

al. [2] who suggested that the ionic conductivity in YSZ is lowered due to the bigger ionic

radius of Y3+ (0.90 Å) as compared to Zr4+ (0.72 Å) resulting in smaller space for the O ions

to move. Since the ionic conductivity also depends on the number ‘n’ of charge carriers

present in the material, therefore based on this, one can postulate that by further increasing

O vacancy concentration in OVSZ, most probably the ionic conductivity will further increase,

as there are no dopants.

Fig. 8. 2: Arrhenius plot of OVSZ thin film containing 16 and 3 at.% Vo. the data reported by

Kosaki et al for 15nm thick YSZ film is also plotted for comparison.

131

This is not the case as shown in Fig. 8. 2. Data reveals that, when the O vacancy concentration

is increased from 3 at.% to 16 at.%, the ionic conductivity is not significantly modified but

its higher than YSZ. The ionic conductivity in case of 16 at.% O vacancy is found to be equal

to ~0.9 S.cm-1 at 650 °C and rises to ~3 S.cm-1 at 725 °C. This behavior in ionic conductivity

for the OVSZ thin film containing 16 at.% O vacancies as compared to 3 at.% O vacancy

concentration film could be explained in terms; i) lower film density of 16 at.% O vacancy

film as compared to 3 at.% O vacancy film as the number of Zr atoms for 16at ,% was found

lower (5.42 E+17/cm2) as compared to 3 at.% (7.5 E+17/cm2) by RBS measurements. ii) by

lattice distortion, induced by the larger number of O vacancies. Such lattice distortion around

O vacancy in pure zirconia has already been reported in several studies [7–10]. Our quantum

– chemistry based calculations results also show that the cubic cell is slightly distorted as the

number of oxygen vacancies is increased as shown in Fig. 8. 3.

Fig. 8. 3: Zirconia cell after relaxation with (a) 0 at.% O vacancy (b) 10 at.% O vacancy, blue atoms

represent Zr atoms and red atoms represent O atoms.

132

Since for high ionic mobility, the lattice has to be highly symmetric [5,6,11], the

incorporation of a larger number of O vacancies (e.g. 16 at.%) in the crystal most probably

leads to a pronounced alteration of the lattice. This situation ultimately results to a relatively

lower ionic mobility of the charge carriers and thus, to a lower ionic conductivity compared

to what is expected for such heavily doped zirconia films. Similar finding was also reported

by Kimpton et. al. [12] who suggested that excessive distortion of the oxygen ion conduction

path by large size ionic dopants is responsible for the decline in ionic conductivity. Further,

the activation energies (Ea) of 10 nm films containing 16 and 3 at.% O vacancies were also

calculated from the slope of the linear fit and were found to be 1.28 and 1.59 eV respectively

(Fig. 8. 2). This puzzling result could be (partially) explained by the fact that having higher

amount of O vacancies means smaller jump to diffuse from position A to B. While in case of

lower amount of O vacancies, the diffusion path from A to B can be longer, so ion would

need more energy.

The influence of the thickness on the ionic conductivity of OVSZ thin films containing 16

and 3 at.% O vacancies is shown in Fig. 8. 4 (a) and (b). It is observed in Fig. 8. 4 (a) and (b)

that the ionic conductivity of OVSZ films increases with the decrease in film thickness. At

725 °C, the ionic conductivity of 100, 50, 20 and 10 nm thick films containing 3 at.% O

vacancies was found to equal ~ 0.04, 0.05, 0.1 and 7.4 S.cm-1, respectively. The ionic

conductivity of 100, 50, 20 and 10 nm thick films containing 16 at.% O vacancies follows

the same trend and was found to be respectively equal to ~ 0.02, 0.06, 0.2 and 3 S.cm-1 at 725

°C. Such an impact of the film thickness on the ionic conductivity is known for oxide ion

conductors and is attributed to the enhancement of ionic mobility due to interfacial

effects[1,6,11,13]. These effects are divided into two categories; (i) the formation of space

charge layer (SCL) at the interface. The SCL is a region of continuum electric charge

133

occurring in dielectrics, where the mobile charge carrier diffuses away by the influence of an

electric field leaving a depletion region of, for example, ionized species such as oxygen

vacancies. Such depletion is electrostatically balanced by accumulation of opposite sign

carriers [11,14]. (ii) The Lattice Strain, induced by slightly or largely mismatched interface

between the substrate and the ion conducting film itself [11]. The degree of disorder at the

interface plays an important role in determining which of these factor is dominant because it

directly controls the formation of defects at the interface, e.g., at the simplest level of analysis,

higher the disorder higher the concentration of the segregated defects at the interface [11].

This higher disorder at the interface caused by lattice strain or by space charge region will

lead to several orders of magnitude greater diffusivity as compared to that of the lattice

[1,6,11,13]. Indeed, it has been shown by Kosacki et al. [1] that such interfacial effects get

more pronounced for films having a thickness lesser than 60 nm. However, we observe a

pronounced increase in the ionic conductivity when thickness of the OVSZ films is

134

Fig. 8.4: Influence of film thickness on the ionic conductivity for a) the 3 at.% vacancy

concentration and b) the 16 at.% vacancy concentration.

135

Fig. 8. 5: Logarithm of the ionic conductivity as function of thickness of films containing 16 and 3

at.% O vacancies, at 725 °C.

decreased from 20 nm down to 10 nm, as shown in Fig. 8. 5. The increase in ionic

conductivity for the 16 at.% O vacancy film is 2 fold as the thickness is reduced from 50 to

20 nm but a dramatic increase (14 fold) is observed when the thickness decreases further

from 20 nm down to 10 nm. Similarly, for the 3 at.% O vacancy film, a colossal (73 fold)

increase is observed when the film thickness is decreased from 20 nm to 10 nm. This increase

in ionic conductivity indicates that the interfacial effects are pronounced for the 10 nm-thick

films.

136

Fig. 8. 6: Influence of interfacial strain on the ionic conductivity of films containing 16 at.% O

vacancies and having various thickness.

To separate the effect related to the space charge layer and the interfacial lattice strain, films

of thickness equal to 50, 20, and 10 nm containing 16 at.% O vacancies were deposited on

MgO (Zirconia - MgO lattice mismatch equals 18 %) and NGO substrates (Zirconia - NGO

lattice mismatch equals 25 %). Ion conductivity data are displayed in Fig. 8. 6 reveal that

films characterized by a thickness of 50 and 20 nm deposited on NGO exhibit a lower

conductivity as compared to the same films deposited on MgO. However, for the 10 nm thick

films, it is observed that the films deposited on NGO (25% lattice mismatch) exhibit a higher

ionic conductivity as compared to one deposited on MgO (18% lattice mismatch). This leads

to the conclusion that the strong enhancement of the ionic conductivity for the 10 nm thick

137

films is more marked by interfacial lattice strain. Above 10 nm, most probably, the lattice

strain relaxes and lead to more bulk (lattice) like diffusion.

Finally, the four-probe analysis performed on the as deposited samples at room temperature

revealed films do not exhibit any electronic conductivity at all i.e. are highly insulating (a

critical requirement for the electrolyte to be used in solid oxide fuel cells).

8.3 Conclusion

Ionic conductivity measurements of OVSZ thin films containing 16 and 3 at.% O vacancies

(1 at.% of O Vac = 3.2 mol.% of Yttrium in YSZ) have been carried out and reveal that OVSZ

thin films are good ion conductors. Ionic conductivity is in the range of reported values for

YSZ. Then increase of O vacancy concentration from 3 at% to 16 at% does not improve the

ion conductivity as first expected. The incorporation of larger amounts of vacancy leads to

lattice distortions which hinders ion mobility. However, a colossal increase in the ionic

conductivity is observed when the film thickness is decreased to 10 nm. This nanoscale effect

seems to be mainly originated from lattice strain, induced by depositing the OVSZ film onto

the right substrate (polished NdGaO3 (100) single crystals). Such an enhancement in the ionic

conductivity of the OVSZ thin films might offer new opportunities for ion conductors.

Moreover, it should be noted that the OVSZ films have a high resistivity as measured by

four-point probe method. This combined with the high ionic conductivity exhibited by OVSZ

makes this material a promising alternate of YSZ to be used in solid oxide fuel cells.

138

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highly textured YSZ thin films, Solid State Ionics. 176 (2005) 1319–1326. doi:10.1016/j.ssi.2005.02.021.

[2] R. Pornprasertsuk, P. Ramanarayanan, C.B. Musgrave, F.B. Prinz, Predicting ionic conductivity of solid oxide

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139

9. Optical properties of OVSZ thin

films

Zirconia is a wide band gap material (bang gap energy = 5.1 - 6 eV[1–5] ) characterized by a

high refractive index (n = 2.1 - 2.2[6–8] ). These properties make zirconia, and zirconia –

based materials, useful candidates for high k dielectric applications, [9–11]. Most of the

zirconia properties are exploited by stabilizing its high temperature tetragonal and cubic

phases with the help of doping, e.g. with Y3+ cations, as previously shown (see chapter 2).

Beside the use of dopants, one can stabilize high temperature c-phase of zirconia at room

temperature by incorporating O vacancies in the zirconia lattice i.e. by growing under-

stoichiometric zirconia (ZrO2-x) thin films, as demonstrated in the chapter 5. In such oxygen

vacancy stabilized zirconia (OVSZ) films, not all the Zr-d electrons are accommodated with

O atoms. This situation leads to the promotion of Zr-labeled states into the bandgap of

zirconia. It is therefore expected that the optical properties of the materials might be modified,

such as in allowing new optical transitions in the visible range. In this last section, we present

the first results on the optical characterization of Oxygen Vacancy Stabilized cubic Zirconia

thin films. In particular, we study experimentally and theoretically how oxygen vacancy

incorporation influences the bandgap and thus optical properties of OVSZ thin films.

140


In order to investigate the optical properties of zirconia thin films, two sets of samples having

thickness 100±10 nm were prepared on Si (100) and on borosilicate glass substrates. Thin

film samples were synthesized at 10 mTorr and with a discharge current of 200 mA. Samples

were prepared by working inside the transition zone as well as out of it, using dc-RMS (details

can be found in chapter 4). Each of these sets is comprised of 3 samples, 2 of which are c-

ZrO2-x samples with oxygen vacancy concentration equal to either 3 at.% or 16 at.% and one

samples is pure ZrO2 (i.e. no vacancies incorporated). In this last case, as previously reported,

the X-ray diffractograms shows only reflection emanating from the monoclinic phase.

For the optical characterization of such deposited samples, spectrophotometry and

photoluminescence (PL) measurements were carried out at room temperature at the Institut

Jean Lamour, in Nancy (Dr. David Horwat and Dr. Hervé Rinnert). In case of

photoluminescence, a Xe lamp was used as excitation source and an edge filter was used as

to cut the excitation wavelength above 325 nm. During the characterization, excitation

wavelengths used are 270 nm (4.6 eV), 280 nm (4.4 eV), 290 nm (4.3 eV), 300 nm (4.1 eV)

and 310 nm (4 eV). Detection was made possible using a CCD camera and, in order to

compare the PL spectra, PL intensities are normalized to the excitation power.

Beside these experiments, data were extracted from the DFT calculations previously

performed to gain a better insight on the experimental results. Details of such calculation can

be found in Chapter 4.


The impact of various amounts of O vacancy incorporated in the zirconia unit cell on the

bandgap, as derived from the DFT calculations, is shown in Fig. 9. 1. At first sight it can be

141

seen the unit cell with 0 O vacancy shows no state in the bandgap. While the DOS of 1, 2, 3

and 10 at.% O vacancy show that, as soon as we incorporate O vacancies in the unit cell, new

states appear in the bandgap. The number of states increases with the increase in O vacancy

concentration. It is also observed that the Fermi level shifts towards the conduction band with

the incorporation of O vacancies. These results underline that these states, located in the

bandgap, are occupied.

Fig. 9. 1: DOS of zirconia unit cells containing 0, 1, 2 3 and 10 at.% O vacancies. Dashed red line

shows the fermi level.

Further analysis of the DOS of the 3 at.% OVSZ film is shown in Fig. 9. 2. It reveals that

these band gap states are not coming from the O atoms but from the Zr atoms. Moreover, it

is also observed that the states appearing in the bandgap are localized near the O vacancy and

belong to Zr atoms (see Fig. 9. 2a, DOS of Selected Zr atoms located close to the O vacancy).

DFT calculations suggest that the formation of an isolated O vacancy in pure zirconia lead to

142

a situation where the O-p band can no longer hold all the Zr-d electrons. As a result, Zr states

accommodating the excess electron appear in the bandgap.

One thing regarding the calculated bandgap has to be noted; its energy is about 3.7 eV.

Experimental values reported for the zirconia bandgap range between 5.1 and 6 eV [1–5].

This discrepancy in the bandgap calculated theoretically and experimentally is well

known[12] and arises due to the PBE functional used in our calculations.

Fig. 9. 2: DOS of cubic zirconia containing 3 at.% O vacancy, calculated using DFT. DOS of (a) Zr

atoms located close to O vacancy, (b) all Zr atoms in the cell, (c) all O atoms, (d) of all atoms.

PL spectra of the deposited films containing 0, 3, and 16 at.% of oxygen vacancies are

reported in Fig. 9 .3 (a), (b) and (c), respectively. PL spectra of 0 at.% oxygen vacancy film

show no emission peak (the very small bump at ~320 nm is due to the parasitic light

emanating from the setup), while the spectra related to the films with 3 and 16 at.% oxygen

vacancy concentrations exhibit two emission peaks. The first peak is centered at 388 nm (3.2

eV) and the second is centered at 488 nm (2.5 eV). The presence of these two emission peaks

143

in the visible range for these two OVSZ films highlight that photoluminescence results from

O vacancy incorporation in the zirconia lattice.

The PL spectra of 3 at.% O vacancy film shows that at low excitation energy (4 eV), emission

intensity of 388 peak is high while the intensity of the 488 nm peak is low. Further, an

evolution of reverse trend is observed with the increase in excitation energy for these two

peaks (Fig. 9. 3b). Since the excitation energy in the present study is smaller than the zirconia

bandgap, therefore, there is no possibility to excite electrons from the valence band (VB)

directly to the conduction band (CB). However, as we discussed, DFT calculations show that

the incorporation of O vacancies leads to the formation of energy states in the bandgap to

accommodate the excess electrons (Fig. 9. 1. Electron from those new states could be excited

to the levels above the CB, leaving holes behind. Those excited electrons decay non-

radiatively to the CB from where they decay back into the holes and lead to the emission of

388 nm (3.2) and 488 nm (2.5) peaks, Fig. 9. 4. However, the DOS of the film doped with 3

at.% O vacancy (Fig. 9. 1) shows that the new states appearing in the bandgap lie at energies

of 1.8, 1.4, and 1.2 eV from the CB. These energies do not correspond to the observed

fluorescence signal wavelengths. As it was discussed earlier, the bandgap calculated here,

using PBE functional, is lower than the reported experimental bandgap energies. To correct

such calculated 3.7 eV bandgap, one way is to lift the CB to the experimental value i.e. to 5.1

eV. By doing so, new values corresponding to the states appearing in the bandgap from CB

become 3.2, 2.8, and 2.6 eV. After this adjustment, one can find the 3.2 and 2.6 eV energy

values. These values fit well with the emission wavelengths hence confirming that the origin

of fluorescence lie in these states appearing in the bandgap as a result of incorporation of O

vacancies. A tentative schematic representation of electron excitation upon light absorption

during the PL experiments, and involving the energy states located in the band

144

Fig. 9. 3: PL spectra of films containing 0 (a), 3 (b) and 16 (c) at.% O vacancy.

145

Fig. 9. 4: Schematic representation of the possible transitions.

gap, is presented in Fig. 9. 4. Further, the decrease in the emission intensity of 388 nm peak

and increase in the emission intensity of 488 nm peak with the increase in excitation energy

for 3 at.% O vacancy film suggest that most probably there is a transfer of energy between

these two peaks.

Fig. 9. 3c shows the PL spectra of the 16 at.% oxygen vacancy film. Those spectra reveal that

the intensity of emission peak at 488 nm increases with the increase in excitation energy,

while the intensity of peak at 388 nm is very weak and excitation energy has no influence on

it. The origin of both peaks is the same i.e. the states appearing in the bandgap as a result of

O vacancy incorporation, as discussed earlier. However the weaker intensity of 388 nm peak

in case of 16 at.% O vacancy and as compared to 3 at.% O vacancy film is probably due to

the quenching of the PL phenomenon by the larger amount of O vacancies present in the

zirconia lattice. It is known for O vacancies that they can trap the excitation energy through

non-radiative transitions and quench the emission [13]. We believe that it’s the reason why

we observe a decrease in the intensity of the peak at 388nm with the increase in O vacancy

concentration.

146

Fig. 9. 5: Transmittance spectra of films containing 16, 3 and 0 at.% oxygen vacancy.

The transmittance spectra of films containing 0, 3 and 16 at.% O vacancy films are shown in

Fig. 9. 5. It is observed, below 270 nm, the samples do not transmit any light due to the glass

substrate which is not transparent to the UV. While in the visible and IR range, films show a

high transmittance above 80 %. Spectra also show small bumps around 900 nm which

correspond to the change of grating and detector. The other small bump around 1050 nm,

probably also coming from the glass substrate. In general, the transmittance data show that

films are transparent in visible-IR range.

9.3 Conclusion

In conclusion, DFT calculations show that incorporation of O vacancy in the zirconia lattice

lead to the formation of occupied energy states in the bandgap. Photoluminescence data

reveal no emission peak in case of 0 at.% O vacancy film while two emission peaks centered

at 388 nm and 488 nm in films containing 3 and 16 at.% O vacancies. This suggests the origin

of 388 and 488 nm peak is the formation of new states in the band gap as a result of O vacancy

147

incorporation. However, having large amount of O vacancies (16 at.%) concentration in the

zirconia lattice leads to the quenching of 388 nm peak. Transmittance data show that the

deposited films are transparent in visible-IR range. Optical data presented here is the very

preliminary data on OVSZ and may subject to more thorough investigation.

148

References [1] K.J. Patel, M.S. Desai, C.J. Panchal, The influence of substrate temperature on the structure, morphology, and

optical properties of ZrO2 thin films prepared by e-beam evaporation, Adv. Mat. Lett. 3 (2012) 410–414.

doi:10.5185/amlett.2012.6364.


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[4] M.M. Larijani, E. Hasani, V. Fathollahi, S. Safa, Thermally oxidized zirconium nanostructured films grown on Si

substrates, Cryst. Res. Technol. 47 (2012) 443–448. doi:10.1002/crat.201100381.

[5] P. Gao, L. Meng, M. Dos Santos, V. Teixeira, M. Andritschky, Characterisation of ZrO2 films prepared by rf

reactive sputtering at different O2 concentrations in the sputtering gases, Vacuum. 56 (2000) 143–148.

http://www.sciencedirect.com/science/article/pii/S0042207X99001992 (accessed October 13, 2014).

[6] S. Zhao, F. Ma, Z. Song, K. Xu, Thickness-dependent structural and optical properties of sputter deposited ZrO2

films, Opt. Mater. (Amst). 30 (2008) 910–915. doi:10.1016/j.optmat.2007.04.001.

[7] S. Ben Amor, B. Rogier, G. Baud, M. Jacquet, M. Nardin, Characterization of zirconia films deposited by r.f.

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37–42.

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[10] J.C. Garcia, L.M.R. Scolfaro, A.T. Lino, V.N. Freire, G.A. Farias, C.C. Silva, H.W.L. Alves, S.C.P. Rodrigues,

E.F. Da Silva, Structural, electronic, and optical properties of ZrO2 from ab initio calculations, J. Appl. Phys. 100

(2006). doi:10.1063/1.2386967.

[11] M.G. Krishna, K.N. Rao, S. Mohan, Structural and optical properties of zirconia thin films, Thin Solid Films. 194

(1990) 690–695.

[12] H. Jiang, R.I. Gomez-Abal, P. Rinke, M. Scheffler, Electronic band structure of zirconia and hafnia polymorphs

from the GW perspective, Phys. Rev. B - Condens. Matter Mater. Phys. 81 (2010) 1–9.

doi:10.1103/PhysRevB.81.085119.

[13] Y. Cong, B. Li, S. Yue, D. Fan, X. Wang, Effect of Oxygen Vacancy on Phase Transition and Photoluminescence

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13974–13978. doi:10.1021/jp8103497.

149

Outlook

Tailoring the crystal structure of any material is an important parameter for controlling its

properties. In this regard, for several decades, various approaches and deposition methods

have been used to understand and control the phase formation of zirconia bulk materials and

thin film coatings; in particular, to stabilize zirconia high temperature cubic phase at room

temperature. For zirconia, phase transition as a function of temperature can lead to

mechanical failure of the material, which in turn, prevent its use in engineering applications.

In the chapters 4 and 5 of this doctoral thesis work, quantum-chemistry based calculations

and optimized reactive magnetron sputter deposition method have been combined to assess

the influence of oxygen vacancy incorporation on the phase formation of zirconia thin films.

According to the calculations, performed at the density function theory level, and their

remarkable agreement with the experimental thin film characterization data, it is concluded

that incorporating O vacancies (at least 3 at.%) in the zirconia lattice is the sole mechanism

responsible for stabilizing the zirconia high-temperature c-phase, at room temperature. The

O vacancy containing zirconia thin films were grown experimentally by dc reactive

magnetron sputtering with a tight control over the oxygen flow (the discharge current was

200 mA and total pressure 10 mTorr). The latter was enabled by implementing a feedback

control loop which links the plasma and discharge parameters to the reactive gas flow, and

to the target oxidation state. This first result highlights the importance of materials defect

150

chemistry during film growth by reactive magnetron sputtering. However, in chapter 6 it is

demonstrated, role of materials defect chemistry is only important and govern the zirconia

phase constitution only when the films are grown in moderate conditions i.e. conditions

providing global energy flux < 1000 eV/Zr atom. Above this energy, the role of materials

defect chemistry is apparently partially hindered and the monoclinic phase formations occurs

and thus, a mixed c- and m-phase is obtained. The increased global energy flux per Zr atom

leading to the build of large compressive stresses in the film is proposed to be the responsible

for such behavior. Beside controlling the materials defect chemistry and deposition

conditions, it is also shown in chapter 6 that monoclinic crystals can also appear when thicker

films are deposited. TEM analysis allowed to detect m-ZrO2 crystals 850 nm above the film-

substrate interface. This shows one also should pay attention to the film thickness when

growing thicker films even if the film deposition conditions are kept constant during the

whole deposition process. In chapter 7, the thermal stability analysis performed on pure c-

phase oxygen vacancy stabilized zirconia (OVSZ) thin films (i.e. 100 nm thin films grown at

200 mA, 10 mTorr) showed that these films are stable to rather high temperatures approx.

750 °C. The appearance of monoclinic diffraction peak after passing 750 °C is attributed to

two possible mechanisms; i) Oxygen incorporation and/or ii) thermal stresses.

Having now the understanding and the means to incorporate oxygen vacancies and to

stabilize the high temperature cubic phase of zirconia, one can further characterize this

material. In this regard, in chapter 8, lateral ionic conductivity measurements are presented.

The results reveal that the OVSZ films are ion conductors, exhibiting conductivity values

similar to those recorded for Yttria Stabilized Zirconia coatings. Furthermore, an ionic

conductivity as high as 7.4 S.cm-1 is obtained at 725 °C for 10 nm-thick OVSZ films

containing 3 at.% of vacancies. Such colossal enhancement in the ionic conductivity is

151

attributed to the lattice strain caused by film-substrate lattice mismatch resulting in higher

disorder at the film-substrate interface. Such epitaxial strain could be a key step in the design

of electrolyte exhibiting high ionic conductivity. Four-probe and Hall effect measurement

showed the OVSZ are highly resistive and exhibit no electronic conductivity at room

temperature. The high ionic conductivity and high electronic resistivity exhibited by OVSZ

thin films makes this material a promising alternate of YSZ to be used e.g. in solid oxide fuel

cells (SOFC) and in oxygen sensors. Further, photoluminescence analysis of OVSZ (chapter

9) revealed the presence of 2 emission peaks centered at 388 nm and 488 nm. This result

validates the theoretical calculations data predicting the appearance of energy states in the

band gap upon O vacancy incorporation. However, the presented photoluminescence data is

preliminary and a thorough investigation is needed to figure out the detailed mechanism and

optical transitions involved in these transitions.

As it is mentioned above, the role of materials defect chemistry is only important and govern

the zirconia phase constitution only when the films are grown in moderate conditions i.e.

conditions providing global energy flux < 1000 eV/Zr atom. Since the global energy flux per

Zr atom is the sum of energies brought to the substrate by all plasma species (Ar+, O-, e-, hv,

IR, atoms, etc.), therefore, it would be interesting to develop a strategy to unambiguously

pinpoint the influence of each energetic plasma specie on the zirconia phase constitution. One

way to achieve this could be the use of mass spectrometry in the current deposition conditions

to get the ion energy distribution function (IEDF) to find the energy of each ionized plasma

specie (Ar+, O-) and relate them with the zirconia phase formation. An ion source to see the

influence of energetic ions on the phase formation of growing film could also be used in this

regard. In similar way, an IR source could also be used separately to inquire the role of IR on

film growth and phase formation.

152

Further, as the increase in film stress as a result of increased global energy flux per Zr atom

is proposed to explain the appearance of monoclinic phase, therefore, it would be interesting

to perform the stress analysis on such deposited films to calculate the critical stress value

leading to the formation of monoclinic phase in these conditions. Moreover, as it is observed

that the increase in film thickness leads to the appearance of zirconia m-phase even if the

films contain 16 at.% O vacancies. Therefore, it would be interesting to figure out the

mechanism responsible for such behavior. One way to understand this mechanism could be

the in-situ stress analysis to assess the stress evolution as the film grows.

Thermal stability analysis of the as deposited cubic zirconia thin films (100 nm thin films

grown at 200 mA, 10 mTorr) showed films are stable up-to 750 °C and after passing this

critical temperature (750 °C) m-phase start to appear. Appearance of m-phase has been

attributed to two possible mechanisms; i) Oxygen incorporation and/or ii) thermal stresses

To unambiguously separate these two mechanisms a more detailed study is needed. This can

be achieved e.g. by measuring the film chemistry before and after the annealing and also by

depositing films on various kinds of substrates characterized with different thermal expansion

coefficients. Further stress analysis should also be performed at room temperature, at critical

temperature (750 °C) and after passing critical temperature to unambiguously identify the

role of thermal stresses.

Besides, understanding of the fundamental mechanisms behind the stabilization of the high

temperature c-phase of zirconia could also pave the way to control the phase constitution of

other metal oxides such as e.g., Bi2O3 in order to promote their high temperature phase to

room temperature. Having stabilized high temperature cubic phase of Bi2O3 (δ-Bi2O3)

without any dopants might also enhance its ionic conductivity. Further, growing under-

stoichiometric CeO2 using the developed method could also help to avoid the use of reduction

153

methods currently implemented to create O vacancies in this material, to enhance its ionic

conductivity.

154

Appendix

Total energy of Zirconia polymorphs as a function of O vacancy, calculated using ab initio

DFT.

Table AI. 1. O vacancies introduced Randomly

No. of Zr

atoms

No. of O

atoms

No. of O

vacancies

Monoclinic

total energy

(eV)

Tetragonal

total energy

(eV)

Cubic total

energy (eV)

32 64 0 -30701.28469 -30701.37687 -30701.13491

32 63 1 -30258.75535 -30259.90955 -30259.28631

32 62 2 -29816.28794 -29820.36984 -29819.04874

32 61 3 -29373.56172 -29378.72436 -29379.91604

32 60 4 -28931.20262 -28936.86533 -28939.41483

32 59 5 -28488.02695 -28495.66794 -28499.19016

32 58 6 -28048.12734 -28057.42449 -28057.97726

32 57 7 -27605.83513 -27614.52185 -27616.83311

32 56 8 -27165.23928 -27174.19326 -27178.35916

32 55 9 -26723.71548 -26735.47371 -26738.11302

32 54 10 -26282.10293 -26291.39565 -26296.25087

Table AI. 2. O vacancies introduced as cluster

No. of Zr

atoms

No. of O

atoms

No. of O

vacancies

Monoclinic

total energy

(eV)

Tetragonal

total energy

(eV)

Cubic total

energy (eV)

32 64 0 -30701.28469 -30701.37687 -30701.13491

32 63 1 -30258.75535 -30259.90955 -30259.28631

32 62 2 -29815.90567 -29820.40569 -29819.03379

32 61 3 -29373.59236 -29379.98953 -29380.09725

32 60 4 -28931.62045 -28937.82173 -28934.86731

32 59 5 -28488.09909 -28498.20545 -28497.62375

32 58 6 -28043.42514 -28057.31826 -28051.31602

32 57 7 -27605.22585 -27616.22044 -27616.79238

155

32 56 8 -27160.72825 -27171.63822 -27165.15831

32 55 9 -26721.98725 -26729.82084 -26736.20141

32 54 10 -26284.63429 -26288.19487 -26294.91857

Table AI. 3. O vacancies introduced apart

No. of Zr

atoms

No. of O

atoms

No. of O

vacancies

Monoclinic

total energy

(eV)

Tetragonal

total energy

(eV)

Cubic total

energy (eV)

32 64 0 -30701.28469 -30701.37687 -30701.13491

32 63 1 -30258.75535 -30259.90955 -30259.28631

32 62 2 -29815.67681 -29819.98039 -29817.53961

32 61 3 -29376.05029 -29378.06347 -29377.39886

32 60 4 -28933.26141 -28934.9256 -28937.66177

32 59 5 -28491.0859 -28496.88414 -28499.94733

32 58 6 -28047.69346 -28055.93488 -28059.69634

32 57 7 -27605.62941 -27612.72958 -27617.79826

32 56 8 -27163.86827 -27174.49147 -27178.58252

32 55 9 -26722.18795 -26730.96115 -26734.71378

32 54 10 -26285.48753 -26292.00887 -26295.94746

156

Fig. AI. 1: Influence of O vacancies on zirconia phase constitution. The energies of each phase are

compared as a function of the concentration of oxygen vacancies. Three ways to distribute the

vacancies in the lattice are presented: (a) random, (b) apart from each others, and (c) clustered.

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