30
ISSN 10683755, Surface Engineering and Applied Electrochemistry, 2012, Vol. 48, No. 6, pp. 491–520. © Allerton Press, Inc., 2012. 491 1 INTRODUCTION The interest in tungsten and its alloys has been driven by its outstanding properties and multiple possi ble applications overviewed by E. Lassner and W.D. Schubert [1]. Tungsten is known having the high est melting point of all metals, T m = 3411 ± 28°C. The macrohardness for polycrystalline tungsten can be up to 650 Vickers with Young’s modulus, E = 390–410 GPa. The roomtemperature ultimate tensile strength (UTS) of polycrystalline tungsten rods is commonly in the range of 580–1470 MPa. Strength, plastic and creep properties of tungsten are of paramount interest for its widespread use as a structural material. The hot strength of pure tungsten is exceeded only by that of rhenium. Alloys based on WReHfC are the strongest man made metallic materials produced at temperatures up to 2700 K. Doped tungsten wires are the most creepresis tant wires, with the exception of monocrystalline tung sten [1]. 1 The article is published in the original. Unfortunately, the electrodeposition of pure tung sten deposits from aqueous tungstate solutions is not viable and possibly hindered by the formation of an oxide layer on the cathode during electrodeposition. That oxide cannot be directly reduced to metallic tungsten, due to a very low overvoltage for hydrogen evolution on tungsten [1]. However, tungsten can be easily codeposited with iron group metals, and publi cations dealing with electrodeposition of tungsten alloys with irongroup metals can be found as early as in the 1930s [2–4]. The first commercial CoW alloy electrodeposition process was reported by Schwartz [5], indicating the role of citrate species as an impor tant factor to achieve the desirable morphology and structure of deposits. It was also shown that the cath ode current efficiency was less than 100%, which sug gested that the hydrogen discharge reaction possibly contributes to the process. Early theoretical studies on W (and Mo) code position with iron group metals were initiated by Brenner [6], Holt and Vaaler [7, 8], Fukushima [9] and Vasko [10]. According to Brenner [11], tungsten Modern Trends in Tungsten Alloys Electrodeposition with Iron Group Metals 1 N. Tsyntsaru a, b , H. Cesiulis c , M. Donten d , J. Sort e , E. Pellicer f , and E. J. PodlahaMurphy g a Katholieke Universiteit Leuven, Department MTM, B3001, Belgium b Institute of Applied Physics, Academy of Sciences of Moldova, Chisinau, MD2028, Moldova c Vilnius University, Dept. Phys. Chem., Naugarduko 24, Vilnius LT03225, Lithuania d Faculty of Chemistry, University of Warsaw Pasteura 1, 02093, Warsaw, Poland e Institucio Catalana de Recerca i Estudis Avançcats (ICREA) and Departament de Fìsica, Universitat Autònoma de Barcelona, E08193, Bellaterra, Spain f Departament de Fìsica, Universitat Autònoma de Barcelona, E08193, Bellaterra, Spain g Department of Chemical Engineering, Northeastern University, Boston, MA 02115, USA email: [email protected] Received January 17, 2012 Abstract—Theoretical and applied studies of tungsten alloys with iron group metals (MeW) are being car ried out worldwide, in the light of their versatile applications. The aim of this paper is to provide an overview of the works on electrodeposition of tungsten alloys with iron group metals, their properties and applications. There are 221 papers reviewing on the following theoretical and practical topics: chemistry of electrolytes used for electrodeposition, codeposition mechanisms, and properties of electrodeposited tungsten alloys. In addition, the formation of W(VI) and iron group metal (Me) complexes (polytungstates and complexes of Me(II) and W(VI)) with citrates and OH is analysed based on the published data and the calculated distri bution of species as a function of pH (ranged from 1 to 10) is provided for solutions with/without citrates. The adduced data are correlated with the compositions of electrodeposited alloys. Various codeposition models of tungsten with iron group metals described in the literature are critically discussed as well. The peculiarities of the structure of tungsten alloys and their thermal stability, mechanical, tribological, and magnetic properties, corrosion performance, their applications in hydrogen electrocatalysis, templateassisted deposition into recesses (aimed to obtain micro and nanostructures) are also reviewed and mapped. DOI: 10.3103/S1068375512060038

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Page 1: Modern Trends in Tungsten Alloys Electrodeposition with ... trends_SEAE491.pdf · cations dealing with electrodeposition of tungsten alloys with irongroup metals can be found as early

ISSN 1068�3755, Surface Engineering and Applied Electrochemistry, 2012, Vol. 48, No. 6, pp. 491–520. © Allerton Press, Inc., 2012.

491

1 INTRODUCTION

The interest in tungsten and its alloys has beendriven by its outstanding properties and multiple possi�ble applications overviewed by E. Lassner andW.D. Schubert [1]. Tungsten is known having the high�est melting point of all metals, Tm = 3411 ± 28°C. Themacrohardness for polycrystalline tungsten can be up to650 Vickers with Young’s modulus, E = 390–410 GPa.The room�temperature ultimate tensile strength (UTS)of polycrystalline tungsten rods is commonly in therange of 580–1470 MPa. Strength, plastic and creepproperties of tungsten are of paramount interest for itswidespread use as a structural material. The hot strengthof pure tungsten is exceeded only by that of rhenium.Alloys based on W�Re�HfC are the strongest man�made metallic materials produced at temperatures up to2700 K. Doped tungsten wires are the most creep�resis�tant wires, with the exception of monocrystalline tung�sten [1].

1 The article is published in the original.

Unfortunately, the electrodeposition of pure tung�sten deposits from aqueous tungstate solutions is notviable and possibly hindered by the formation of anoxide layer on the cathode during electrode�position.That oxide cannot be directly reduced to metallictungsten, due to a very low overvoltage for hydrogenevolution on tungsten [1]. However, tungsten can beeasily codeposited with iron group metals, and publi�cations dealing with electrodeposition of tungstenalloys with iron�group metals can be found as early asin the 1930s [2–4]. The first commercial Co�W alloyelectrodeposition process was reported by Schwartz[5], indicating the role of citrate species as an impor�tant factor to achieve the desirable morphology andstructure of deposits. It was also shown that the cath�ode current efficiency was less than 100%, which sug�gested that the hydrogen discharge reaction possiblycontributes to the process.

Early theoretical studies on W (and Mo) code�position with iron group metals were initiated byBrenner [6], Holt and Vaaler [7, 8], Fukushima [9]and Vasko [10]. According to Brenner [11], tungsten

Modern Trends in Tungsten Alloys Electrodeposition with Iron Group Metals1

N. Tsyntsarua, b, H. Cesiulisc, M. Dontend, J. Sorte, E. Pellicerf, and E. J. Podlaha�Murphyg

aKatholieke Universiteit Leuven, Department MTM, B�3001, BelgiumbInstitute of Applied Physics, Academy of Sciences of Moldova, Chisinau, MD�2028, Moldova

cVilnius University, Dept. Phys. Chem., Naugarduko 24, Vilnius LT�03225, LithuaniadFaculty of Chemistry, University of Warsaw Pasteura 1, 02�093, Warsaw, Poland

eInstitucio � Catalana de Recerca i Estudis Avançcats (ICREA) and Departament de Fìsica, Universitat Autònoma de Barcelona, E�08193, Bellaterra, Spain

fDepartament de Fìsica, Universitat Autònoma de Barcelona, E�08193, Bellaterra, SpaingDepartment of Chemical Engineering, Northeastern University, Boston, MA 02115, USA

e�mail: [email protected] January 17, 2012

Abstract—Theoretical and applied studies of tungsten alloys with iron group metals (Me�W) are being car�ried out worldwide, in the light of their versatile applications. The aim of this paper is to provide an overviewof the works on electrodeposition of tungsten alloys with iron group metals, their properties and applications.There are 221 papers reviewing on the following theoretical and practical topics: chemistry of electrolytesused for electrodeposition, codeposition mechanisms, and properties of electrodeposited tungsten alloys. Inaddition, the formation of W(VI) and iron group metal (Me) complexes (polytungstates and complexes ofMe(II) and W(VI)) with citrates and OH– is analysed based on the published data and the calculated distri�bution of species as a function of pH (ranged from 1 to 10) is provided for solutions with/without citrates. Theadduced data are correlated with the compositions of electrodeposited alloys. Various codeposition models oftungsten with iron group metals described in the literature are critically discussed as well. The peculiarities ofthe structure of tungsten alloys and their thermal stability, mechanical, tribological, and magnetic properties,corrosion performance, their applications in hydrogen electrocatalysis, template�assisted deposition intorecesses (aimed to obtain micro� and nanostructures) are also reviewed and mapped.

DOI: 10.3103/S1068375512060038

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SURFACE ENGINEERING AND APPLIED ELECTROCHEMISTRY Vol. 48 No. 6 2012

TSYNTSARU et al.

codeposition with iron group metals—Me (i.e., Ni,Fe, Co) from aqueous electrolytes is called “inducedcodeposition”. Citrate complexes were identified byBobletsky and Jordan [12, 13] and later used to explainthe species involved in the electrodeposition of W(Mo) alloys by Holt and co�workers [14, 15]. Othertypes of electrolytes have also been found useful fortungsten alloys electrodeposition, namely: malic, tar�taric, gluconic, glucaric [16–19], acetic [20], glycine[21–24] and pyrophosphates [24].

Theoretical and practical studies of tungsten alloyswith iron group metals (Me–W) continue to be carriedout worldwide, in the light of versatile applications ofthese alloys due to their specific mechanical, tribolog�ical, magnetic, electrical and electro�erosion proper�ties as well as corrosion resistance. Figure 1 capturesthis growing interest in this area as the number of sci�entific papers published in the top�rated journals.Me–W alloys may compete even with ceramics andgraphite by virtue of their high thermal resistance.Furthermore, tungsten alloys with a nanocrystallinemicrostructure make up mechanically hard coatingsthat have favorable wear and anti�corrosion proper�ties, which make them extremely attractive for replac�ing hard Cr coatings in many applications whereaggressive conditions exist [18]. That includes pro�cesses where lubrication or dry friction conditions takeplace, for example in hydraulic equipment, as well asthose where a lubricant may be difficult to maintain,for ex., in aerospace drives and splines where low fric�tion in dry conditions over a limited period of flightmay be a valuable asset [18]. Tungsten alloys can alsobe used as a barrier layer for ultra large�scale integra�tion devices (ULSI) [25] and micro�electro�mechan�ical systems (MEMS) structures [26, 27], and as cata�lysts for clean hydrogen generation from water split�ting [28]. It is worth noting that Me–W shows also atunable magnetic behavior, ranging from soft or semi�hard ferromagnetic properties (for moderate W con�

tents) to non�ferromagnetic, when the W contentexceeds 25–30 at % [29].

The aim of the present paper is to provide an over�view of the researches on electrodeposition of tungstenalloys with iron group metals, their properties andapplications.

CHEMISTRY OF ELECTROLYTES USED FOR W�ALLOYS ELECTRODEPOSITION

Understanding the chemistry of the electrolytesused for alloy electrodeposition is important to unveilthe factors controlling the performance of the platingbath and the peculiarities of codeposition of alloy con�stituents.

Any practical use of electrolytes containing onlyiron�group metal salts (i.e., “simple” baths with e.g.,

sulphates, chlorides) and without ligands islimited due to the formation of sparingly solublehydroxides and tungstates of corresponding iron�group metals (Me–Co, Ni and Fe); these hydroxidesprecipitate at pH > 7.5–7.7 [30], and the iron�grouptungstates precipitate at pH > 2–3 [31–33].

A number of different W(VI) polynuclear speciesare present in water when tungstates are dissolveddepending on the pH values:

—at pH > 7.8 ion is the predominant spe�

cies in solution, whereas is the main speciesat pH < 7.8, and it is in a protonated form at pH < 5.7

[34]. Also, (paratungstate�A) and (paratungstate�B) can exist [35];

—at pH < 5 equilibria are slowly established and

and are dominating [35–36];

—at pH 4 protonates into which with time forms the stable metatungstate ion

[37];

—at pH 1 the main tungsten species in solution isthe hydrated tungstic acid WO3 ⋅ 2H2O [38].

In Fig. 2 we provide the calculated mole fractionsof W(VI)�containing species vs. pH in a solution con�taining 0.1 M Na2WO4. The calculation is based on theavailable equilibrium constants including tungstic acid

and tungstates [36, 39]. The amount of “free” is reduced sufficiently at pH < 5 and at pH 2–3 the

concentration of decreases down to 10–10 M

due to the formation of species. However, ifan iron�group metal is added as a sulfate salt, the ion�

pairing (or ion�associates) Me2+ : will domi�nate, which is not the case in chloride or nitrate solu�tions. This reduces the concentration of “free” Me2+�ions down to 10–3 M or even below in a 0.2 M MeSO4

WO42–

WO42–

W12O4212–

W7O246– H2W12O42

10–

W12O4110– HW6O21

5–

HW6O215– H3W6O21

3–

H2 W3O10( )4[ ]6–

WO42–

WO42–

W12O396–

SO42–

19900

1992 1994 1996 1998 2000 2002 2004 2006 2008 2010

510152025303540

Years of publication

Number of papers

Fig. 1. Number of papers on theoretical and practicalaspects of W alloys electrodeposition published in theissues referred by “ISI Web of Knowledge” database (WEBof Science® and Science Citation Index ExpandedTM).

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SURFACE ENGINEERING AND APPLIED ELECTROCHEMISTRY Vol. 48 No. 6 2012

MODERN TRENDS IN TUNGSTEN ALLOYS ELECTRODEPOSITION 493

solution [40]. Therefore, solutions containingNa2WO4 and MeSO4 are stable at pH < 3 because ofthe MeWO4 solubility product Ksp: e.g. Ksp = 10–13 forFeWO4 [41]. Such solutions can be used to obtain W�containing alloys, e.g., Fe–W [42–43] or Co–P–W[44]. When pH rises to 3–6, the concentration of

“free” ions increases up to 10–7–10–5 M, whichis high enough to exceed the solubility product, Ksp, ofMeWO4, and therefore insoluble MeWO4 are formedand precipitate in the absence of complexing ligands.

Complexing is used as a method to approach theelectrode potentials of different metals by reducing theconcentration of “free”Me2+ ions and increasing thesolubility of Me(II) salts in baths containing OH–,

and [44, 45]. Several common electro�lytes are known and used for the electrodeposition oftungsten alloys with iron�group metals but the numberof complexing agents used is limited (citrates, pyro�phosphates [23, 24], tartrates [16, 17], gluconates[18, 19]). Usually, a certain formulation can beadapted for the codeposition of Co, Ni and Fe alloyswith W, as well as for the electrodeposition of ternaryalloys [40, 46–49] by simple exchange of iron groupmetals salts. When citrates or pyrophosphates areusing as complexing agents, the concentrations ofCo2+ and Ni2+ in solution decrease 104–106 timescompared to those in ligand�free solutions [40, 50].The added citrate modifies the distribution of tung�sten�containing species, and various complexes of thetype [(WO4)p(HCit)qHr]

(2p + 3q – r)– are formed. Thestability constants of eight tungstate�citrate complexesare provided in [51].

In Fig. 3 we show the calculated mole fractions oftungsten species in an equimolar tungstate�citratesolution vs. pH. Under these conditions at low pH,

only one polytungstate complex, , is formed insufficiently high amounts. In the entire range of

WO42–

WO42– PO4

3–

W12O396–

pH values, different tungstate�citrate complexes areformed and the main ones are the following:

—[(WO4)1(HCit)1H3]2–, [(WO4)2(HCit)2H5]

5–

and [(WO4)2(HCit)1H4]3– at pH 2–4;

—[(WO4)2(HCit)2H6]2– and [(WO4)1(HCit)1H2]

3–

at pH 4.5–7.5;

—[(WO4)1(HCit)1H]4– at pH 7–9.

Only at pH > 9, the species becomes pre�dominant. A similar distribution of tungsten�contain�ing species should remain also in the citrates�contain�ing solutions for tungsten alloys electrode�position,where the concentrations of iron�group metals are rel�atively low in comparison with those of citrates; andsimultaneously, the concentrations of citrates arecomparable with those of tungstate, e.g., for Ni–W[46, 47, 52–56], Co–W [46, 57–59], Fe–W [46, 60, 61],and ternary alloys [47, 49].

The content of Me and W in the alloy depends onthe ratio of the partial current densities (PCD) fortungsten and Me. Moreover, the PCD for tungsten cancorrelate with the PCD of the iron�group metal [62].

When the concentration in solution is higherthan the Me salt concentration, the PCD for W is thehighest under such conditions and, therefore, depositswith relatively high W contents are obtained, irrespec�tive of the total current applied, e.g., at least 20–25 at %W for Ni–W, and 30–35 at % W for Co–W and Fe–Walloys. However, the Faradaic efficiency (FE) is rela�tively low and in most cases does not exceed 10–15%.

In order to increase both the FE and depositionrate, solutions containing higher concentrations of theiron�group metals salts (0.1–0.2 M) are used whereas

WO42–

WO42–

10

3 5 7 9

0.2

0.4

0.6

0.8

1.0

pH

Fraction of W in species12*HW12O41

10–

12*HW12O396– WO4

2–

6*HW6O215–

Fig. 2. Mole fraction of tungsten species in a 0.1 MNa2WO4 solution.

10

3 5 7 9

0.25

0.50

0.75

1.00

pH

Fraction of W in species

(WO4)(HCit)H4–

(WO4)2(HCit)H43–

(WO4)2(HCit)2H46–

(WO4)(HCit)H32–

(WO4)(HCit)H41–

(WO4)2(HCit)2H55–

(WO4)(HCit)H23–

(WO4)(HCit)2H62–

WO42–

12*W12O396–

Fig. 3. Mole fraction of tungsten species in solution withequimolar quantities of Na2WO4 and citrate.

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the concentration of citrates is kept at the level of 0.3–0.5 M [44, 45, 63–66]. This brings about an increasein the FE up to 30% for Fe–W, up to 33–50% for Ni–W,and 60% for Co–W. Under these conditions, the dis�tribution of species in solution differs from the onepresented in Fig. 3. Citrate complexes both with tung�state and iron group metals are leaving a sufficient frac�tion of “free” Me2+ ions (at pH < 3–4) or ionic pairs

“Me2+ : ” (at pH < 7–8) in solution. As an exam�ple, the calculated mole fractions of tungsten(VI) andCo(II) species are provided in Figs. 4a–4c. For theNi–W case, the calculated distributions of W(VI) andNi(II) are similar to those for Co(II). In these solu�tions low soluble tungstates of iron�group metals arenot formed within the entire range of the working pH

due to either low concentrations at a low pH orlow Me2+ ions concentrations at a higher pH. The

SO42–

WO42–

increase in concentration with the solution pH(Fig. 4c) correlates with the increase in W content inthe alloys deposited at the same fixed overall current den�sity at pH ranged from at least 5 to 8.5 [45, 64, 67–69].Indeed, the dependence of W content in alloys vs. pHat the same total current density has the maximum atpH 8–8.5, that cannot be predicted based on the

increase of [ ] with pH (Fig. 3) at pH<6.5 or the

plateau in [ ] at pH > 6.5 (Fig. 4c). It can be sup�

posed that ions might form other complexes ata higher pH. Thus, the codeposition data were ana�lyzed by assuming the formation of a mixed complex[(Ni)(WO4)(Cit)(H)]2– in solution [70, 71]. In fact, anelectrolyte that contains citrate complexes of cobaltand tungsten may include the formation of com�pounds with enhanced molecular mass, probably

WO42–

WO42–

WO42–

WO42–

1 2 3 4 5 6 7 8 9 10

0.25

0.50

0.75

1.00

01 2 3 4 5 6 7 8 9 10

0.25

0.50

0.75

1.00

0

1 2 3 4 5 6 7 8 9 101E–10

(WO4)(HCitr)H32–

(WO4)(HCitr)H4–

WO42–

(WO4)(HCitr)H23–

(WO4)2(HCitr)H43–

W12O396–

(WO4)2(HCitr)2H4

(WO4)(HCitr)2H6

(WO4)2(HCitr)2H55–

(WO4)(HCitr)H41–

pH

Fraction of W(VI) in species(a)

pH

Fraction of Co(II) in species

Co2+

:SO42–

Co2+

CoH2Cit

CoH3Citr+

CoHCitr–

(b)

1E–08

1E–06

0.0001

0.01

1

pH

Fraction of Co(II) and W(VI) in species

WO42–

Co2

(c)

CoOH+

Fig. 4. Mole fraction of tungsten (a) and cobalt (b) species, as well as probable charge�transfer species (c) for Co–W electrodepo�sition. Composition of the solution: 0.2 M CoSO4, 0.2 M Na2WO4, 0.3 M citrates.

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SURFACE ENGINEERING AND APPLIED ELECTROCHEMISTRY Vol. 48 No. 6 2012

MODERN TRENDS IN TUNGSTEN ALLOYS ELECTRODEPOSITION 495

polynuclear heterometallic complexes containing aniron�group metal (e.g. Co) and W(VI), as discussed in[72]. However, the stability constants of these hetero�nuclear complexes are not presented in this reference.In addition, the slow adsorption stages occurring onthe surface during both pure iron�group metal [73]and W alloys electrodeposition [74] hinders an accu�rate prediction of tungsten content as a functionof pH.

Ammonia (NH3) is frequently added to platingbaths to fine�tune the pH within the range of 7.0–9.0and to increase the FE, especially in the case of Ni andFe alloys electrodeposition. In the case of Co alloys,the distribution of Co(II) complexes in solution withor without NH3 is almost the same, so that little effectis impaired on the FE of Co–W [40, 50]. Nevertheless,ammonia plays a key role in the baths, namely as abuffering agent. The composition of W�containingelectrodeposited alloys is very sensitive to the solutionpH [67, 69, 75, 76]. In the case of W alloy electrodepo�

sition, the electroreduction of �ion and hydro�gen evolution reactions are associated with increasingpH at the electrode surface, which yields the electro�lyte alkalization in the boundary layer near the elec�trode that can sufficiently increase or decrease the Wcontent in the alloys, depending on the initial pH. Aconstant alloy composition along the entire thicknessof electrodeposits can be achieved if the OH– ionsreleased during the electroreduction of tungstate fur�ther participate in the fast acid�base equilibrium reac�tions with hydrogen ions, keeping the pH in the nearelectrode layer almost constant. In other words, theelectrolyte must possess a high buffer capacity; other�wise, compositionally graded deposits would beexpected. The buffer capacity as a function of pH in thecitrate�ammonia system was analyzed in [44, 45, 67]. Itwas shown that at pH < 7.5 the buffer capacity is con�trolled by the concentration of citrates. Therefore, forthe electrodeposition of W alloys from acidic solutionsit is unnecessary to use ammonium salts as bufferingadditives. Usually the electrodeposition of W alloys iscarried out at elevated temperatures. For the ammoniaconcentration of 0.8 M, the minimum buffer capacityat 70°C is not as deep as at room temperature.

The maximum in buffer capacity shifts to lower pHvalues if the electrolyte temperature is increased. Forexample, deposition at near room temperature and atpH of 8.5 would not be a good choice, whilst at 70°Cthe deposition works fine. In order to reduce possibleevaporation of ammonia from the electrolytes duringlong�term electrolysis at elevated temperatures andthus to keep the alloy deposition rate stable and offsetthe increase in pH values due to the reduction of

and the hydrogen evolution, it is recommendedto perform electrodeposition at pH values as low aspossible (7.5–8.0). In this case, a sufficient fraction ofthe ammonia added to the solutions transforms into

the non�volatile �ionic form. Moreover, because

WO42–

WO42–

NH4+

of better buffering properties of solutions at these pHvalues, a higher concentration (1.5 M) of ammonia isrecommended as well [45, 67]. The electrodepositionfrom aqueous baths is generally carried out in non�hermetic cells. In this case, however, significantamounts of ammonia could evaporate during long�term electrodeposition at elevated temperatures,which results in a decrease of both pH and depositionrate. In order to avoid the decrease in pH due toammonia evaporation, Na2CO3 can be introducedadditionally to the electrolyte to maintain the desiredpH. In the absence of Na2CO3, the amount of W in theNi–W alloys is more sensitive to pH than to the cur�rent density, and there is a large difference in depositcomposition at pH between 8.5 and 10. Therefore,adding Na2CO3 to the bath reduces the effect of bothpH and the current density on W amount in the Ni–Walloys [45].

As the presence of NH3 is not critical for Co–Walloys electrodeposition, boric acid is added to thesolution for buffering purposes instead of ammonia,so this solution does not contain volatile components[63]. In this solution, the buffer capacity remainsalmost constant at the level of 0.2–0.4 mol/L in therange of pH 6 to 11. At 70°C, the buffer capacity in anarrower pH range, i.e. from 7 to 9 is even higherthan or the same as that was obtained at lower tem�peratures [64].

Generally, the chemical composition of the elec�trodeposited alloy depends on the ratio of the PCD foralloy’s constituents. However, the value of the PCD isinfluenced by a number of factors such as chemistry ofsolution, total current density applied, temperature,hydrodynamics, etc. Therefore, at a given pH, theincrease of overall current density may result either in adecrease [68, 77], increase [18, 19, 64] in the tungstencontent in the alloys, or it can remain almost constant[65, 67].

CODEPOSITION MECHANISMS

The composition of an alloy resulting from an elec�trodeposition process is a key parameter for the designof requisite material properties, and is a direct conse�quence of the ratio of metal species reduction reactionrates (i.e., partial current densities) that occur. Under�standing the phenomena that determine the PCD isthus essential to be able to predict and control the dep�osition. The induced codeposition behavior of tung�sten alloy electrodeposits from aqueous media intro�duce a further complexity in that both kinetic andtransport mechanisms can be coupled. In this section,the kinetic mechanisms and the coupled nature ofmass transport effects are discussed.

One of the earliest reports to describe the kineticmechanism of tungsten codeposition with iron�groupelement was made by Holt and Vaaler [8]. A “catalyticreduction” theory was proposed to explain the reductionof aqueous tungstate solutions in the presence of the zero

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SURFACE ENGINEERING AND APPLIED ELECTROCHEMISTRY Vol. 48 No. 6 2012

TSYNTSARU et al.

valence state of an iron�group element, Me (i.e., Ni, Coand Fe). The following two cathode reactions are sug�gested as being essential to the reduction process:

Me2+ + 2e– Me, (1)

(2)

The iron�group element reduction occurs first,which then catalyzes tungsten deposition, until thesurface is replete with tungsten, arresting the tungstendeposition. This proceeds until a new layer of the iron�group element is formed. The observation of lami�nated structures germinated this concept, although asstated by Younes and Gileadi [78], non�laminated,solid solution alloys of tungsten with iron�group ele�ments can also be readily electrodeposited. That sem�inal work [8] was also reviewed by Vasko [10]. Hepointed out that while metals of the iron�group areactive catalysts, it is also known that W can co�depositon other, very inactive metals (e.g., Mn, Sb, Cu andSn) but to a small extent. Another important implica�tion of the assumption that the solid phase of an iron�group metal acts as a catalyst, is that the amount of Wthat can be theoretically codeposited into an alloy isrestricted to an inherent upper limit. A review of earlierworks in [10] provides these upper W wt % limits: 53 wt %with Ni, 65 wt % with Co, and 80 wt % with Fe.

While Holt and Vaaler [8] did recognize that tung�sten codeposition occurs with Faradaic efficienciesless than 100%, due to the hydrogen evolution sidereaction, they did not place a direct dependence ofhydrogen on the tungsten mechanism. In contrast,Clark and Lietzke [79], emphasizing the role of hydro�gen, presented a different mechanism where, first,there was a deposition of a partially reduced tungstatefilm on the cathode, followed by the catalytic reduc�tion of this film by hydrogen in the presence of freshlydeposited iron, cobalt, or nickel. Without the inducingelement, the tungstate ion is just partially reduced, anddue to its low overpotentials to hydrogen evolution,cannot reduce further. They also stated that the react�ing species were those that were complexed and didnot rule out iron, cobalt or nickel tungstate species.

Iron�group metal tungstate complexes that are pre�cursors to tungsten alloys is a concept introduced byYounes and Gileadi [78], Younes�Metzler et al. [70],and reviewed by Eliaz and Gileadi in [80]. The idea ofimplicating the mixed�metal soluble complexes as aprerequisite for iron�group and tungsten codepositionwas motivated by the observation that the partial cur�rent density of Ni depends on the amount of tungstatein the electrolyte and vice versa. In publications onNi–W codeposition, Gileadi et al proposed a complex

composed of Ni2+, and deprotonated citrate.While it was recognized that bimetallic complexes canexist with citrate, none has been identified with aniron�group metal and tungsten. Soluble complexedspecies as the means for codeposition of Co and W hasalso been suggested by Aravinda et al. [81]. They iden�tified a cobalt citrate complex with UV absorption

WO42– 8H+ 6e– W 4H2O.+ + +Me

WO42–

studies, but not a mixed�metal complex, and toexplain shifts in cyclic voltammetry curves, they pos�tulated that alloy deposition is possibly from a cobalt�tungstate complex. A gluconate (gluc) containingelectrolyte for Co–W depostion was examined byWeston et al. [18] with UV spectroscopy and by com�paring spectra with different amounts of added speciesthey deduce that significant quantities of a mixed

metal complex occurs, [Co–gluc– ]. Theyobserved a decrease in the reduction current densitywith an increase in the quantity of the ternary mixed�metal complex and concluded that this one may bemore difficult to reduce than a binary complex. Theyattributed the high W content in the deposit directlyfrom the reduction of a ternary complex, consistentwith Younes and Gileadi [78], and Eliaz and Gileadi[80] points of view. In addition, Belevskii et al. [82]detected a heteropolynuclear Co–W citrate complexthrough the evaluation of gel chromatography.

However, the formation of various complexes doesnot explain the variety in dependencies of W contentin the alloys on the electrodeposition conditions. TheVasko film model reviewed in [10, 11] is an alternateinterpretation. For instance, when employing Ni–W,the formation of a nickel hydroxyl species, NiOH+,NiOHads aids the adsorption of tungstate species:

Ni2+ + OH– NiOH+, (3)

NiOH+ + e– NiOHads, (4)

n ( )ads. (5)

A heteropolytungstate film is formed with a certaincomposition that is not defined, [HpNiqWrOs]w givingthe following overall, two�step reaction:

(6)

(7)

The ensuing high resistance in the cathode surfacecreates a high negative potential at the metal (alloy)–film interface. The heteropolytungstate film is graduallyreduced to the alloy and at the same time new layers ofthe film are continuously formed. Since films were notalways observed during deposition, this model haslacked popularity. However, recently, Jušk nas et al.[84] reported a transient in the Ni–W composition dur�ing film growth that was mapped by in situ X�ray dif�fraction (XRD). They state that the Vasko model fitsthese best, and has renewed interest in this concept.

The induced codeposition behavior of tungstenalloy reduction has often been considered similar tomolybdenum alloy reduction and analogies of mecha�nisms from molybdenum studies warrants discussion.Early Ni–Mo electrodeposition was also studied bysome of the same authors. Ernst et al. [14], and Ernstand Holt [15] proposed that the induced elements, Moor W, formed a mixed oxide intermediate at the elec�trode surface which was subsequently reduced byhydrogen to solid metal. This approach stems from the

WO42–

WO42– WfOg

y–

NiOH+ 3H3O WO42– HpNiqWrO5[ ]w,+ +

HpNiqWrO5[ ]w ne– Ni–W mO2– nOH–.+ + +

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observation that generally as more Mo (or W) is code�posited into an alloy the hydrogen side reaction tendsto be higher, or in other words, the current efficiency islower. The mechanism was expanded by Fukushimaet al. [85] assuming that hydrogen was adsorbed ontothe codeposited iron�group element and that theinducing element was reduced at these sites. Thisassumes that the solid iron�group metal acts as a cata�lyst in reducing Mo and that a theoretical upper limitexists for Mo codeposited in an alloy. More recently,this group [86] suggested that the electrodepositionbehavior of W�iron�group metal alloys can be explainedby the reduction mechanism with atomic hydrogensimilar to the induced codeposition of Mo. An earlierwork showed that the adsorbed hydrogen onto thereduced iron�group metal, H(Me), induces the reduc�tion of the molybdenum oxide intermediate [18].

Me2+ + 2e– Me, (8)

H+ + e– H(Me). (9)

+ (6 – n)e– + Mon+ oxide(hydroxide), (10)

Mon+oxide(hydroxide) + nH(Me) Mo(Me). (11)

Assuming a tungsten oxide intermediate with a +4valence state, four hydrogen atoms are needed toreduce a W(IV) oxide to a W atom and every Ni atomcan adsorb 0.6 hydrogen atoms, which is equal to thenumber of unpaired 3d electrons in Ni. Thus, whenthe atomic hydrogen bonds to the Ni atom, this con�tributes to the reduction of the W(IV) intermediate atthe ratio of W atoms to Ni is 0.6/4 (32 wt %). A higheramount of W(IV) oxide intermediate forms whenmore tungstate is in the solution and to reduce this theNi atoms have to contribute to the reduction severaltimes. Oue et al. [86] found by x�ray photo�electronspectroscopy (XPS) that unreduced tungsten oxideaccompanied solid W at high tungstate electrolyteconcentrations and that the Ni(II) reduction rate wassimultaneously inhibited, and in turn decreased thereduction rate of tungsten. The deposited Ni–W alloyshad two distinct regions, depending on the relation�ship between the alloy composition and the currentefficiency: a region of a low W content in the deposits,where the partial current efficiency for W depositionincreases with the W content, and a region of a highW content in the deposits, where the partial currentefficiency of W decreases with an increase in W content.

Chassaing et al. [87] suggested that the iron�group ionwas responsible for the formation of the intermediate, notthe solid state, and that the intermediate was reduced byhydrogen forming solid Ni and Mo, codeposited as analloy. At low overpotentials, where only an oxide wasformed, Auger electron spectroscopy (AES) and energydispersive x�ray spectroscopy (EDS) identified a molyb�denum�nickel oxide intermediate having a 4 : 1 ratio ofNi to Mo. They suggested the following induced codepo�

Me

MoO42–

sition mechanism: when no alloy deposit is formed, thefollowing two reactions occur:

+ 2H2O + 2e– MoO2 + 4OH–, (12)

MoO2 + 4Ni2+ + 8e– MoO2Ni4 (13)

at larger overpotentials, when an alloy deposit isformed, then the overall reaction proceeds as:

(14)

with a mixed�metal intermediate.More recently, Go�mez et al. [88] showed that the

first stage of the alloy formation involved partialreduction of Mo to different oxidation states consis�tent with that reported by Fukushima et al.[86] and,Chassaing et al. [87] In both the Fukushima and Chas�saing models there is an upper limit for Mo (or W)codeposition in an alloy. To date, the highest amountof W codeposited with Ni is 50–67 a/o (76–86 w/o) Win a Ni–W alloy, as reported by Younes and Gileadi[78], and the highest amount of Mo codeposited withNi is 50–74 a/o (62–82 w/o), as reported by Sun andPodlaha [89], although both with impractically lowFaradaic current efficiency.

While it was possible to find deposition conditionswhere an increase in Mo concentration in the depositwas accompanied by a drop in the current efficiencyand thus hydrogen evolution, it was also possible tofind the contrary situation, calling into question thedecisive role of hydrogen [90]. A different modelingapproach was proposed by Podlaha and Landolt [91,92], which explained the induced codeposition of Mo(or W) via a catalysis reaction, assuming the catalystwas a mixed Mo (or W)�iron�group intermediate, butwithout the need for the adsorbed hydrogen. Themixed�metal intermediate could adsorb at the elec�trode surface and was subsequently reduced withoutany controlling feature being dominated by the hydro�gen evolution side reaction. The mechanism was ini�tially reported for Ni–Mo electrodeposition [15] froma citrate electrolyte, and then generalized to othercodepositing iron�group elements (Fe, Co) [92];assuming a complexed nickel ion species with a cit�rate ligand, NiCit–, which was not intended to cap�ture the details of complexation, the mechanism waspresented as:

NiCit– +2e– Ni(s) + Cit–, (15)

(16)

(17)

where Cit– represents [C6H5O7]3–, and while the

NiCit– species is to represent the complexation ofNi(II) with a citrate species, the exact form of the spe�

MoO42–

MoO42– 3Ni2+ 4 n–( )H 8e–+ + +

→MoNi3 4OH– nH+ +

MoO42– NiCit+

–2H2O 2e–+ +

NiCitMoO2[ ]ads– 4OH–

,+

NiCitMoO2[ ]ads– 2H2O 4e–+ +

Mo s( ) NiCit– 4OH–.+ +

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cies will, of course, depend upon the electrolyte com�position and pH. That mechanism should be viewed asbeing cognizant of complexing equilibria without spe�cific details. Similarly, the Ni species could also bereplaced with Co or Fe. The molybdate species in theabove mechanism could be equally replaced withtungstate species to make this model more general. Ifassuming that the iron�group species acts as a catalyst,and that it is reduced independently to the mixed�metal ion intermediate, then no limitation is placed onthe quantity of Mo (or W) that can be codeposited.Go�mez et al. [88] took this mechanism one step fur�ther and suggested that the NiCit– complex speciescatalyzes the molybdenum oxide that is initiallyformed, combining the features of the Chassaing andPodlaha�Landolt models. With this assumption, whenthere is no NiCit– present in the electrolyte, the par�tially reduced molybdenum oxides are predicted.

The Podlaha–Landolt model does exclude theimportance of the side reaction and confirmes that,indeed, the adsorbed intermediates of hydrogen canalso compete with the mixed�metal intermediates,indirectly influencing the reaction rate. For example,when molybdenum or tungsten alloys are tested forcatalyzing hydrogen evolution, the Tafel slopes can berelated to conventional hydrogen evolution kinetics;Jaksic [93] noted that on NiMo alloys hydrogen evo�

lution follows the Volmer�Heyrovsky mechanism inbase electrolytes, with a high surface coverage by theadsorbed hydrogen, according to:

H2O +e– Hads + OH–, (18)

Hads + H2O + e– H2 + OH–. (19)

If assuming a similar mechanism occurring duringdeposition of the alloy, then the adsorbed hydrogenmay be significant enough to affect the available freesurface sites for the Mo (or W) and the iron�groupintermediates. In a similar manner, the deposition ofthe iron�group elements can also occur with coverageof an adsorbed intermediate. For example, Ni, Fe andCo elemental electrodeposits have been widely mod�eled via an adsorption intermediate [94–100]:

Me(II) + e– [Me(I)]ads, (20)

[Me(I)]ads + e– Me (s). (21)

and, hence, can also alter the free electrode surfacearea available for the Mo (or W) reduction reaction.

The utility of the Podlaha–Landolt model is bestshown in the coupling of the kinetic and transportregimes of the molybdenum reduction reaction (seeFig. 5). When the concentrations of the inducing irongroup ions (e.g., Ni) in the solution are very low com�pared to those of molybdate (or tungstate), then theMo (or W) reduction would be limited by the amount

Inducedcodeposition

Inducedcodeposition

Elemental deposits Alloy deposits

Uncoupled

iW, i2

i1, i2

CW(VI) > C2

C1 � C2

#2

E

iNi

(a)

(b)

(c)E

E

EE

E

iW, i2

#1

#1#2

#2

#2

#2#2

CW(VI) � C2

CW(VI) � C2

iW iW

iW

iW

iNi

i2Ni

δ

δ

δ

Fig. 5. Comparison of PCD in a kinetically, uncoupled (a) versus induced codeposition (b) behavior alloy reduction when onespecies, #2, is reduced under mass transport control. A change in the boundary layer thickness alters only species #2 in an uncou�pled alloy codeposition, while it affects both species #2 and the inducing element, e.g., W.

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MODERN TRENDS IN TUNGSTEN ALLOYS ELECTRODEPOSITION 499

of available inducing ions. In addition, species that arein low concentrations quickly reach their mass trans�port limit (i.e., limiting current density). Thus, inuncoupled reactions, if one metal ion species is inexcess, and another species is, in contrast, having avery low concentration in the electrolyte, it is possibleto electrodeposit in the regime where the species inexcess exhibit a kinetic control and the other exercisesa mass transport control during codeposition. In thiscase, an increase in the electrolyte agitation willincrease the amount of the minor species in thedeposit due to its increased limiting current density(Fig. 5a). However, the induced codeposition behaviordoes not mirror this expectation. If the minor speciesin the electrolyte are any of the inducing elements(e.g., Ni, Fe or Co) and if it is under mass transportcontrol, then with an increase in the electrode or elec�trolyte agitation, its partial current density willincrease and so will the reluctant metal (Mo or W)reaction rate, keeping the deposit concentation rela�tively insensitive to agitation (Fig. 5b). Or said anotherway, the inducing element appears to be transport con�trolled even though its concentration in the electrolyteis in excess and is unexpectedly influenced by suchtransport factors as changes in the hydrodynamicenvironment. The opposite scenario, when the reluc�tant element, such as the tunsgtate species is the minor

component, it behaves as expected under transportcontrol with an increase in its PCD, as though it isapparently decoupled as long as an iron�group ele�ment is codepositing (Fig. 5c).

To capture this effect, a steady�state model thatintegrates both transport and kinetics is needed. Anexample is given in Table 1, that follows the secondmodel presented by Podlaha and Landolt [92], wherethe fractional adsorption of the reacting complexediron�group metal, θML, was included along with theadsorption of the mixed�metal molybdate intermedi�ate, θMo–ML. The reduction reaction rates for theinducing element was assumed to be governed by theTafel kinetics, and a Nernstian boundary layerapproach was adopted where diffusion dominateswithin the boundary layer, and the boundary layer size,δ, is established by convection induced by the hydro�dynamic environment near the electrode surface fromelectrode or electrolyte agitation.

Figure 6 illustrates the transport influence when Niions are in low concentration compared to molybdatecodeposited onto a rotating cylinder electrode,CNi(II) � CMo(VI), hence the boundary layer can be

defined by δ = 99.62 where di isthe electrode diameter in cm, ν is the kinematic vis�

di–0.4ν0.344Dlim

0.356S–0.7,

Table 1. Summary of the Podlaha�Landolt model of induced codeposition [92]

Boundary conditions at the electrode Governing equationsBoundary conditions at the boundary layer

thickness, δ

∇ ⋅ NML = 0

assuming a Nernstian approach:

for each species

NML

iML 1,

F����������

iMo 1,

4F����������

iMo 2,

F����������–+= Cj Cj

b=

NMoO4

2�

iMo 1,

4F����������= ∇ N

MoO42�⋅ 0=

iML 1,

F���������� kc ML 1,, CML

s1 θML– θMo–ML–( )e

αc ML 1,,

F

RT������������������E

⎩ ⎭⎨ ⎬⎧ ⎫

–= N DdCdx�����–=

iML 2,

F���������� kc ML 2,,– θMLe

αc ML 2,,

F

RT������������������E

⎩ ⎭⎨ ⎬⎧ ⎫

=

iMo 2,

F���������� kc Mo 2,, CML

sC

MoO42�

s1 θML– θMo–ML–( )θMLe

αc Mo 1,,

2F

RT���������������������E

⎩ ⎭⎨ ⎬⎧ ⎫

–=

iMo 2,

F���������� kc Mo 2,, θMo–MLe

αc Mo 2,,

4F

RT���������������������E

⎩ ⎭⎨ ⎬⎧ ⎫

–=

iside

F������� kc side, CML

s1 θML– θMo–ML–( )θMLe

αc side,

F

RT����������������E

⎩ ⎭⎨ ⎬⎧ ⎫

–=

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cosity in cm2 s–1, Dlim is the diffusion coefficient of thespecies that is under or near its limiting current densityin cm2 s–1 and S is the electrode rotation rate in rpm. ThisNi–Mo example is similar to the scenario in Fig. 5b. Atransport influence is expected in the Ni partial cur�rent density due to its low concentration in the elec�trolyte and high polarization of the electrode. If thiswas an uncoupled reaction, with an excess of molyb�date in the electrolyte there should be no change of theMo partial current density with the electrode rotationrate, but there indeed is. Thus, the Mo partial currentdensity experiences a change in deposition rate due tothe coupled nature of the Ni reduction reaction. Thesolid lines in Fig. 6a, 6b are the modeling results of theequations given in Table 1, assuming that the side reac�tion and the Ni codeposition reactions have interme�diates with very low surface coverage.

If tungsten reduction occurs primarily through aternary, mixed�metal complex of some type, [Me�

ligand� ], as presented by Younes and Gileadi[78], and Younes�Metzler et al. [70], Weston [18] andBelevskii [82], as opposed to an adsorbed intermedi�ate, the apparent transport effect of the W PCD, whenthe concentration of the iron�group metal in the elec�trolyte is much less than the added tungstate species,CMe(II) � CW(VI), would still manifest, as in Fig. 5b.Hence, the W PCD would be limited by the availableinducing ions to form the complex, and both metalPCD would mirror each other. This effect can be seenin the results of those reported by Younes and Gileadi[78] for the Ni–W codeposition system in citrateammonia electrolytes. In the opposite scenario, ifthere was an excess of the inducing, iron�group ele�

WO42–

ment(s), CMe(II) � CW(VI), the formation of a ternary Wcomplex would be limited by the amount of availableW. Since the iron�group metal can be reduced inde�pendently, eqns (20) and (21), the PCD behaviorwould be represented by that in Fig. 5c, giving anapparent decoupling of the W PCD with the iron�group element. To simulate essentially the samebehavior as the Podlaha–Landolt model in Table 1,one needs to modify the governing equations in thebulk and at the electrode surface to account for thecomplexing equilibria.

Differentiating between an adsorption model or acomplexing species model for the induced code�posi�tion of W (or Mo) is not straightforward with steadystate polarization data. Electrochemical ac impedancespectroscopy does give an indication that adsorbed spe�cies are present, through the observation of an addi�tional capacitive or inductive feature during Ni–Mo[92, 97], Ni–Fe [92], and Co–W [74] codeposition,which provides credibility to the adsorption modelapproach as adopted in the Podlaha–Landolt model.Impedance results of Chassaing et al. [87] were inter�preted to be representative of a porous surface layer onNi–Mo, which is more in line with a film model suchas the one proposed by Vasko [10].

A limitation of the model in Table 1 is that it doesnot account for observable decreases of the PCD atvery high overpotentials. Beltowaska�Lehman [98]observed that there is a decrease for both Ni and MoPCD at applied potentials more negative than –1.3 Vvs SCE. It is most pronounced with an increase inmolybdate concentration, and attributed to the for�mation of molybdenum oxides. Analogously, Obra�dovic et al. [99] observed a similar decrease in the total

0.70

0.8 0.9 1.0 1.1 1.2 1.3 1.4

1

2

3

4

5

–E, V vs NHE

Ni partial current density, mA/cm2

0.70

0.8 0.9 1.0 1.1 1.2 1.3 1.4

2

4

6

8

10

–E, V vs NHE

Mo partial current density, mA/cm2

Rotation Rate, rpm30070013002000model

Rotation Rate, rpm30070013002000model

(a) (b)

Fig. 6. Partial current density of Ni and Mo, under the influence of a kinetic�transport control, showing an increase in the (a) Nipartial current density with electrode rotation rate and an induced rotation rate dependence, or apparent transport influence of

the (b) Mo partial current density, when is in excess in the electrolyte; reproduced from ref. [91], courtesy of the Elec�

trochemical Society.

MoO42–

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MODERN TRENDS IN TUNGSTEN ALLOYS ELECTRODEPOSITION 501

current density at high applied overpotentials in theNi–W alloy system. In addition, the simulation of thePCD does not capture the region near the open�cir�cuit potential and hence the thermodynamics of thesystem. Kroger [100] outlined changes in the activitiesdescribed in the Nernst equation for systems thatundergo abnormal deposition, including inducedcodeposition.

PROPERTIES OF TUNGSTEN ALLOYS WITH IRON GROUP METALS

Structure and Thermal Stability

The galvanic technique gives a wide opportunity toproduce alloys of specific texture and nano� or amor�phous internal ordering. Amorphous and nanocrystal�line alloys have some specific mechanical properties.The grain size of electrodeposits depends on the elec�trodeposition technique used [101, 102], the chemicalcomposition and pH of the electrolyte [47, 103] thecurrent density [45], as well as on the substrate [104].However, a specific internal structure of those types ofhard and wear resistant materials is often combinedwith their severe brittleness.

The most critical factor determining the structureof electrodeposited tungsten alloys is the quantitativecomposition of the alloy. It was proved that by opti�mizing the direct current (DC) plating conditionstungsten alloys with Ni, Co or Fe result in the forma�tion of thermo�dynamically stable intermetalic com�pounds—Ni4W, Co3W and Fe2W [48]. Those types ofelectrodeposited alloys are usually fully amorphous,hard and smooth films [47]. Investigations on thestructure of electroplated Ni�W alloy performed byDonten et. al [46] indicated that this alloy has fibrous(columnar) structure and contains ultra�thin fibers(ca. 5 nm in diameter) oriented perpendicularly to thesurface of the substrate. Single fibers aggregate andform broader columns when the thickness of thedeposit extends to 1–2 micrometers. Such a phenom�enon is well visible for Ni–W and Co–W alloys and lesspronounced for Fe–W ones. This type of structure hasimplications on the physical properties of these materi�als. It was proved that the tungsten amorphous alloyexhibits strong magnetic anisotropy [105], breaks mucheasier in the direction normal to a substrate [106] and,because of its strong texture, probably has differenthardness when measured in different directions.

When the alloys feature less amount of tungsten,they deposit in the form of a solid solution. In thiscase, tungsten is dissolved in the iron�group metal lat�tice. The material retains the structure of the basemetal—Ni and Co are of the fcc type (Fm3m) and Fehas a bcc lattice (Im3m). The structures of deposits areoften nanocrystalline with increased lattice parame�ters compared to pure iron group metals [86]. Usually,16–18 at % of W in Ni or Co are the critical upper val�ues for the formation of solid solutions. Changes in the

alloys structure at these mentioned compositions,associated with the occurrence of a fully amorphousstructure, have been reported in [107–108]. Thedeposits have also a strong (111) texture and exhibit anincrease of (111) interplanar distance associated withan expansion of the lattice parameters, that is in theconformity with Vegard’s law [47, 86]. In this case,increasing the tungsten content in the deposited alloyalso decreases the crystal grain size, ranged from a sub�micrometers in materials containing several at % of Wto 5–10 nm, when the amount of tungsten exceeds15 at % [44].

Increasing the tungsten content beyond 18 at %disturbs the solid solution structure of the electro�plated tungsten alloys. Huang et al. [109] showed thata Ni–W electrodeposit with tungsten content at leastof about 18 at % may consist of the Ni4W intermetallicphase. Krolikowski et. al. [110] found out that the cor�rosion behavior of a nickel�tungsten alloy is signifi�cantly changing at ca. 18 at % of W. At a higher tung�sten content, the structure of intermetallic com�pounds probably becomes highly deformed. This phenomenonis described at large for Ni–W alloys. Sridhar et. al[111] also defined conditions for electroformation ofNi4W. For tungsten content from about 20 to ca.35 at %, electrodeposited alloys are fully amorphous[112–113]. Also Co–W and Fe–W alloys with a hightungsten content reveal amorphous (for Co–W) [68]or nanocrystalline structures, with the nanocrystal sizesmaller than 5 nm for Co–W [64] and for Fe–W[65, 114].

Tungsten alloys deposited in a pulse current modealso show an amorphous structure in the range of 18–35 at % of W; when compared to DC plated alloys,they turn to be more uniform and probably have a lesspronounced columnar structure [58, 64]. Any furtherincrease of the tungsten content in pulse electrode�posited alloys leads to the formation of other inter�metallic compounds. Such deposits containing morethan 35 at % of W are described in a number of papersoriginating from Gileadi’s research group. In general,while any chemical composition consisting of around50 at % of W may contain a NiW phase, an increase ofthe tungsten content to about 66 at % leads to theNiW2 phase. Younes et al. [71, 115] considered the for�mation of an orthorhombic NiW phase by XRD mea�surements, when the concentration of W in the depositexceeded 40 at %. It was argued that formation ofequal amounts of Ni and W in this case may beregarded as evidence, albeit circumstantial, for theexistence of a tertiary nickel�tungstate�citrate com�plex. Later, Younes et al. also reported the first platingof a body�centered tetragonal NiW2 phase [70]. As itwas shown on XRD patterns presented in the afore�mentioned papers, alloys containing more than 40 at %of tungsten were not amorphous. It is worth noticingthat alloys with the tungsten content exciding 40 at %were obtained from the ammonia�free plating bath.

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The influence of ammonia and alkaline cations on theplating process and the structure of deposited alloyswere considered by Zakharov et al. [116].

It should be mentioned that the composition oftungsten�containing iron�group�metal alloys andtheir structure for some substrate materials can changewith the thickness of the deposit (distance from thesubstrate) [46, 104, 117]. The similarity in crystallo�graphic structure of the substrate and the selectedcomponents of the electroplated alloy (e.g. coppersubstrate and nickel from electrodeposited Ni–Walloy) can locally change the composition and struc�ture of the deposits. However, these changes, visible onthe cross�section of galvanic coverage, are limited tofilms not thicker than ca. 1 micrometer. Occurrence ofthis interaction between an electrodeposited layer andits substrate does not usually change the amorphousstructure of the deposited layer but significantlyimproves its adhesion to the substrate [104].

Heat treatment can also change the structure andphysical properties of tungsten alloys. Electrodepos�ited amorphous and nanocrystalline materials are usu�ally thermodynamically unstable and tend to increasetheir grain size [48]; they often recrystallize when sub�jected to thermal treatment. The most severe changesin the internal structure are observed in amorphousalloys. Their metastable glass form is changed to apolycrystalline state of a lower potential energy [118].Yamasaki et al. [119] noticed a slow, systematicincrease of size in nanocrystals of Ni4W during anneal�ing of Ni75W25 electrodeposited alloy. Up to tempera�ture of 600°C, the increase of the crystal size was veryslow. Longer heating of the material at 500–600°Cresulted in the formation of ordered grains not largerthan 12 nm in diameter. Above this temperature, theeffect of recrystallization was significantly accelerated.After keeping the Ni–W alloy at 700°C for 24 h, themean crystal size increased to ca. 150 nm. Differentialthermal analysis (DSC) indicated two exothermicphase transformations of nickel�tungsten alloy at tem�peratures of 610°C and 830°C. Those two signals wererecognized as the formation of NiW and NiW2 inter�metallic compounds. This indicates that the amor�phous structure of deposited alloys slowly transformsto nanocrystalline Ni4W during annealing at moderatetemperatures. In this process, the hardness of the coat�ing significantly increases from a value of 770 HV (as�deposited amorphous alloy) to 1450 HV (annealednanocrystalline alloy with the grain size 12 nm).

Similar DSC experiments, supported by the obser�vation of modifications in the magnetic properties ofthe electrodeposited Ni80W20 alloy, indicated phasetransformation of the amorphous material appearingat a lower temperature range (400 to 500°C) [120]. Inthis case, two well�shaped DSC signals at tempera�tures of 431°C (0.34 J/mol) and 445°C (0.45 J/mol)were explained by the formation of the crystallineNi4W phase. Further increase of temperature, up to

800°C, indicated serious changes in the alloy structureand its magnetic susceptibility at temperature close to650°C. Changes in hardness during the annealing pro�cess probably caused by structural modifications wereobserved during heat treatment of binary cobalt�tung�sten and ternary cobalt�iron�tungsten alloys [121]. Inthis case, the highest hardness in alloys containing ca25 at % of W was noticed after annealing at 600°C.

Exposing Ni–W and Co–W alloys to much highertemperatures leads not only to an increase of the crys�tal size of the solid solution and formation of interme�tallic phases, but also to decomposition of the alloywith separation of carbides and metallic tungsten[114, 122]. As reported by Tsyntsaru et al. [64], ther�mal stability of Co–W alloys depends on the contentof tungsten in the deposit. While the deposits contain�ing 24.5 at % of W began the recrystallization processat 400°C, the alloys with higher ca. 35 at % W retainedtheir amorphous structure to almost up to 600°C.Similar transformation was also observed for elec�trodeposited Fe–W alloys. It was found out thatrecrystallization of alloys containing 23–30 at % of Wstarts at temperatures above 400°C, but the nanocrys�talline structure of the deposits does not change up toca. 800°C. Elevation of the heating temperature up to1000°C resulted in decomposition of the alloy to formFe, Fe2W and FeWO4 crystalline phases [114]. The lat�ter could be caused by recrystallization of oxidespresent on the surface and in cracks of the electrode�posits.

Results of investigations performed on ternaryNi⎯Fe–W and Fe–Co–W alloys indicated that signif�icant, temperature�induced changes of those alloysoccur at temperatures above 500°C and 700°C,respectively. The recrystallization of the amor�phous/nanocrystalline Ni–Fe–W alloy [123] con�taining 10–14 at % of W started at temperatures500°C–600°C and led to the formation of α�FeNiphase, Ni4W and Fe7W6 intermetallic compounds.Because annealing of the alloys was performed in air, alarge amount of material was oxidized to NiWO4 andFeWO4. Investigations on the thermal stability of Fe–Co–W alloy containing 50 at % of Fe, 33 at % of Coand 17 at % of W [124] revealed no changes in thestructure at temperatures up to 600°C. Conversely, sig�nificant changes were observed for deposits annealedat temperatures above 800°C. In this case, the forma�tion of nanocrystalline α�Fe phase was observed inXRD patterns and differential thermal analyses (DTA)plots. At the same temperature, the hardness of thesealloys reached its maximum (ca. 1300 HV). Furtherincrease of temperature resulted in fast recrystalliza�tion, with the formation of α�Fe, cubic Co andFe6W6C phases.

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Mechanical and Tribological Properties

Tungsten alloys with iron group metals possessattractive mechanical and tribological properties [18,52–54, 56, 59, 63, 77, 106, 114, 121, 124–130]. A pos�sibility of them to substitute conventional hardchrome plating from electrolytes containing Cr(VI)due to the toxicity of the hexavalent chromium speciesand volatility of them at high operating temperature ofthe chromium bath is often discussed in literatureamong a number of other applications of these alloys[18, 52, 53, 121, 131]. Industrial parts that are oftenhard chrome�plated include aircraft engines and land�ing gear, oil well equipment, crank shafts, hydrauliccylinders, paper making equipment, molds, stamps,dies, drill bits, and power industry equipment. Thealternatives to hard chromium from aqueous�basedelectrolytes might be [131]:

—chromium coatings obtained from Cr(III) basedsolutions;

—nickel�base coatings (electroplated and electro�less), alloys and composites;

—cobalt�base electroplates;—plasma electrolytic oxidation.In this section, hardness and wear resistance of

electrodeposited tungsten alloys are mainly addressedto evaluate recent developments in investigation ofnickel� and cobalt�base alloys with tungsten. Frictionalso is an important property, but considering and esti�mating the available data, it should be taken intoaccount that it is partially a property of a sliding systemthat depends on a variety of interrelated mechanical,chemical, physical, and surface properties of slidingmaterials and surfaces [132]. Mechanical behaviordepends on the deposit structure, composition as wellas on the electrodeposition operating conditions. Thefactors influencing the tribological and mechanicalbehavior can be broadly classified into internal andexternal factors.

Internal factors include recrystallisation tempera�ture, crystal structure, grain boundary effects, etc.,which significantly influence the shear strength of amaterial [134].

Recrystallisation temperature:In some cases hardness is evaluated after heat treat�

ment of coatings [124, 130], so direct comparison withdata on as�deposited alloys is not straightforward[63, 114] because the structure changes after heating.

The highest hardness has been obtained for Fe–Walloys (1400 HV [114]), followed by heat�treatedFe⎯Co–W coatings (1140 HV [124]). In a descendingorder, we find Co–W alloys (800–1000 HV [18, 63]).

Nevertheless, this hardness is comparable to thehardness of the electroplated chromium under similarconditions of nanohardness measurement. Lower val�ues were obtained for as�deposited Ni–W coatings[77, 129]. Ni–W alloy electrodeposition has beenunder investigation by many groups of researchers

[52–54, 56, 99, 115, 128, 138–139]. However, somedecrease in hardness after annealing was detected forNi–Fe–W alloy [140]. Using metallic microstructuresas mold inserts for hot embossing and injection mold�ing requires hard and wear resistant materials in orderto achieve an exact replication into polymers andceramics [138]. The hardness and thermal stability ofmold inserts were found to increase significantly withnickel�tungsten alloys. Nickel�tungsten alloys wereemployed for the manufacture of micro�structuredtools due to their excellent mechanical propertiesregarding wear and mechanical durability [138]. Thesetools have found applications in hot embossing andinjection molding processes. As was previously noted[115], the concentration of W in the plated alloy has amajor effect on its mechanical and chemical proper�ties, such as hardness, abrasion resistance, andimproved corrosion resistance at high temperatures.Then, the authors in [139] reported that the micro�hardness of Ni–W deposits was increased with anincrease in the tungsten content in the alloy. Obra�dovic et al. [99] reported that Ni–W alloys exhibitsuch enhanced properties as corrosion and wear�resis�tance, which makes them useful in practical applica�tions.

Crystal structure:Tribological behaviour depends on the crystal

structure. For example, Wang et al. [135] comparednanocrystalline Ni and Co coatings upon fretting atequal grain sizes and microhardness and noticed thatthe wear rate of Co is much lower than that of Ni coat�ings. One of the possible reasons for this behaviorstems from the difference in the crystalline lattice;namely, cobalt has hexagonal close�packed crystal lat�tice, which gives better tribological performance.Nevertheless, the performance of Ni coatings can beimproved by alloying it with tungsten, as has beenaforementioned [52, 54, 56].

Grain size and boundary:

The Hall�Petch relation, H ~ d–1/2, correlates thehardness of a metal, H, with the inverse square rootpower on the mean grain size, d. However, the param�eter d has been used interchangeably with the term“crystallite size” in the literature under review due toestimation of the particles sizes from XRD spectra.Still, as shown by Schuh [56], when particles sizes arerelatively small (~10 nm), then crystallite and grainsizes obtained by means of the XRD and transmissionelectron microscopy (TEM) become close to eachother. Thus, in the present paper the term “grain size”will be kept, because the majority of the publicationson tungsten alloys discussed (see Table 2) is in the“safe region” of less than 50 nm.

In the case of nanocrystalline alloys, at least twomajor factors contribute to hardening of those alloys:in the case of nanocrystalline alloys, there are at leasttwo major factors contributing to hardening of thosealloys: (1) formation of a solid solution (i.e., solid

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solution hardening), and (2) small average grain size.In the case of tungsten alloys (e.g. Ni–W) the harden�ing due to the formation of a solid solution has aweaker effect compared to hardening due to decreasedgrain size [56]. Similar investigations were made onNi–W alloys by Yamasaki [54], for which the averagegrain size was even smaller (2.5–7 nm were attainedwhen the tungsten content was the highest) than thosein [56] (6–10 nm). It was observed that, when thegrain size is decreasing from the micrometer rangedown to the nanometer range ~ 10–15 nm, the Hall�Petch relation is conserved, namely, with decreasingthe grain size the hardness will increase [133]. Whenthe grain size is below 10 nm, the Hall�Petch break�down is observed, and a modulation and slight weak�ening of mechanical properties can be noticed [133].Thus, Schuh in [56] stated that the Hall�Petch break�down is observed for Ni–W coatings with grain sizes

below ~8 nm. This value is lower than that of pure Ni(~14 nm), suggesting that alloying has suppressed theHall�Petch breakdown Nevertheless, the mechanismof this breakdown is not well understood as some of theinterrelated factors can affect the behavior of the coat�ings: structural changes, diffusion creep, the triplejunction (at grain boundaries), increasing dislocations[134]. Thus, the mechanisms of deformation and theproperties of the nanocrystalline materials depend onthe average grain size, but are also influenced by thegrain size distribution and the grain boundary struc�ture (low/high�angle grain boundaries) [56, 134].

External factors like temperature, humidity, inden�tation size effect and contact conditions [132] alsoplay a role and, hence they should be taken into con�sideration when assessing tribological properties, e.g.,depending on the electrodeposition conditions impu�rities as carbon, sulfur or hydrogen can be incorpo�

Table 2. The relationship between structure and hardness for tungsten alloys

Conditions Alloy Concentra�tion at % W

“Grain size”, nm

Hardness, HV Load, g Reference

Citrate�ammonia solution pH 9.2; 60°C

Co–W 2–15 – 450–680 10 [59]

Citrate�gluconate solution pH 6.0; 80°C

Co–W Fe–Co–W 27 – 538–848 25 [121]

Citrate�borate solution pH 6.0; 40°C, pulse deposition

Co–W 0–7 6–50 800–1000 200 [125]

Citrate�ammonia solution pH 7.5 Co–W 10–17 14–19 620 5 [77]

Borate solution pH 3.5–6; 30–60°C

Co–W 13 220–450 600 100 [126]

Citrate�borate solution pH 6.7; 60°C

Co–W 18–22 150–200 600–725 10 [127]

Citrate–ammonia solution pH 8.5, 75°C

Ni–W 8.4–12.7 20–22 500–570 100 [52]

Citrate�ammonia�borate solution pH 8.5; 75–85°C

Ni–W 1–16 5–50 640 100 [53]

Citrate�borate solution pH 8, 80°C Fe–Ni–W 7–28 200–300 550–630 300 [128]

Citrate�ammonia solution 72°C Ni–W 11–13 4–10 611–815 1 [56]

Citrate�ammonia solution pH 7.5; 75°C

Ni–W 18–23 8–10 560–690 200 [54]

Citrate�gluconate solution pH 6; 80°C

Co–W 14–25 5 500–1000 25 [18]

Citrate�ammonia solution pH 7–8, 70°C

Co–W and Fe–W 28 3–42 600–800 1–20 [106]

Citrate�borate solution pH 6.75; 58°C

Co–W 17–23 5–40 500–845 0.2–20 [63]

Citrate�ammonia solution pH 9.5, 65°C

Ni–W 15–30 – 650–800 15 [129]

Citrate�ammonia solution pH 8, 50°C

Ni–Fe–W 9–14 – 450–600 10 [130]

Citrate�ammonia solution pH 8; 70°C

Fe–W 23–30 3–5 700–1400 1–20 [114]

Citrate�ammonia solution pH 8; 80°C

Fe–Co–W – – 827–1141 98 [124]

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MODERN TRENDS IN TUNGSTEN ALLOYS ELECTRODEPOSITION 505

rated into the coating [136] and have an impact oncontact conditions during sliding and on hardness val�ues as well. Additionally, copious hydrogen evolutionfrom the side reaction during electrodeposition causesproblems associated with micro� or nanovoids forma�tion in the deposit, that sometimes yields contradic�tory results [133]. Besides, other parameters such asdifferences in the loads applied, initial roughness anddifferent set�up for measuring mechanical and tribo�logical properties can affect the measured values.Depending on the materials applied, hardness of thesamples is a design�technological factor. The mutualarrangement of material hardness and counterbodymay correspond to either a so�called direct frictionpair (movable sample is harder than the fixed one) or areverse friction pair (movable sample is softer than thefixed one). In higher kinematic friction pairs (line orpoint contact) the combination of materials “hard–hard” is also used.

Thus, the comparison of tribological and mechan�ical properties based on the published data is quite

complicated because of different starting conditions ofalloys deposition. A few studies can be compared toassess the wear rate of different tungsten alloys with thesame set�up e.g. [63, 66, 114]. The Fe–W, Co–W andCo–W–P coatings were investigated in ball�on flatsliding tests against corundum ball counterbody. TheCo–W coating possesses the lowest wear rate, evenhaving the hardness a bit smaller (at 20 mN) than thatof Fe–W or ternary alloys (Figs. 7–9). Noticeably, thegrain sizes and tungsten content are very close to eachother for the investigated alloys (see Table 2), suggest�ing that the alloying metal (Fe, Co) or grain boundarystructure have an influence on the final hardness andtribological behavior. Weston et al. [18] also comparedthe tribological behavior of electrodeposited chro�mium coatings, nanostructured Co and Co–W elec�trodeposited alloys against steel counterbody. Thewear rate of Co–W coatings was minimal in compari�son to the electrodeposited pure cobalt and chromiumcoatings. Even at the maximum load, while the chro�mium coatings were worn out, Co–W underwent lessdamage. The results of that study are suggesting that

100 μm100 μm (b)(a)

Fig. 7. SEM images of coatings after dry sliding under normal load of 2 N: (a) Co–W and (b) Fe–W alloys.

0

2

4

6

8

10

12

14

16

Nanohardness, GPa2 mN20 mN

CoW

FeW

CoW

P

Ele

ctro

lyti

c ch

rom

ium

Fig. 8. Average nanohardness of tungsten alloys in compar�ison with electrolytic chromium and Co–W–P. Graphsare prepared based on the data published in [66].

0

10

20

30

40

Wear rate, 103 μm3/J

Ele

ctro

lyti

c ch

rom

ium

Co�

W�P

TiN Fe�

W

Ti(

Al,

N)

TiB

2 Co�

W

Dia

mon

d

Fig. 9. Wear rate of tungsten alloys in comparison withother hard coatings. Graphs are prepared based on the datapublished in [66].

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Co–W alloys are best suited to replace chromiumcoatings.

The Fe–W alloy undergoes a severe tribo�oxida�tion during dry friction because iron oxidation causesan expansion of the wear track and cracks appear (seeFig. 7). Nevertheless, the tribological performance ofFe–W coatings can be highly improved under lubri�cated conditions [137] or by using ternary alloys suchas Fe–Co–W [121, 124]. It should be mentioned thatthe authors in [124] did not compare binary Co–W orFe–W coatings with ternary Fe–Co–W coatingsunder the same conditions; hence the comparisonwith other publications is a rough generalization.

Certainly all these factors, both internal and exter�nal, are interrelated in between them. The tempera�ture of electrodeposition, pH, electrolyte compositionand current (direct or pulse current or pulse reversecurrent [3, 10]) or potential applied do exert certaininfluence on the overpotential value and partial cur�rent of tungsten reduction and, eventually, on nucle�ation process of the given alloy, thus changing the crys�tal structure or preferential grow of crystals, grain sizeor grain boundaries. Thus, the growing interest inelectrodeposited tungsten alloy coatings and struc�tures, e.g, Co–W alloys [18, 59, 63, 77, 106, 121, 125,127, 131] and Ni–W alloys [52–54, 56, 115, 99, 128,138–139] will be further increasing because of thecombination of mechanical, tribological and magneticproperties that these materials demonstrate in addi�tion to corrosion resistance and catalytical properties.

Corrosion Properties

Today materials for applications are frequentlyselected according to their functional properties. Theelectrodeposition of films and electroforming arepromising methods to be applied in galvanotechnics.Due to a large surface–area�to�volume ratio of thinfilms, as compared to that of bulk materials, the prop�erties of thin films, as a rule, significantly differ fromthe bulk (macro�scaled) materials behavior. Filmcomposition, microstructure and density are verymuch dependent on preparation methods and condi�tions. The variation in corrosion rates and other prop�erties of the surfaces with the fabrication method andstructure are well known [141, 142]. Then, differencesin corrosion properties between cast iron�group met�als and the corresponding electrodeposited metalshave been observed. The values of corrosion potential(open circuit potential—OCP) are more negative forthe electrodeposited metals than for the correspond�ing cast metals. Furthermore, higher corrosion cur�rents are obtained in electrodeposited metals [143].Therefore, the examination of the corrosion proper�ties of the obtained materials is an integral part of themodern materials science and crucial in terms of tech�nological applications. In the meantime, the uniquecombination of tribological, magnetic, electrical andelectroerosion properties of tugsten alloys has trig�

gered an increased potential of these materials inmacro�, micro�, and nanotechnics. Indeed, a funda�mental requirement for materials in microtechnologyis that they show an extremely high corrosion resis�tance. This is due to the continuous decrease in thecomponent size, down to the micrometer range.Under such conditions, the resistance to device failuredue to even a small amount of ionic contaminants maybecome a key issue in terms of lifetime and reliability.

Tungsten (as well as molybdenum) alloys show arelatively low hydrogen overvoltage [144], and corro�sion of tungsten (and molybdenum) alloys with irongroup metals undergoes mainly with hydrogen depo�larization. Moreover, the absolute values of bothhydrogen evolution reaction and anodic reaction haveto occur at the same rate, while leaving the alloy in thecorrosive media. These are important factors govern�ing the relatively negative values of corrosion potentialand the rate of corrosion, which in some cases (e.g.Co–W) correlates well with exchange current densityfor hydrogen evolution [145].

Tungsten alloys with iron group metals hardlytransfer into a passive state in neutral solutions[110, 146, 147]. The shapes of Evans diagrams of allinvestigated freshly electrodeposited alloys reveal anon�passive state [66, 148]. Even after exposing thesematerials for 15 days in open air, the corrosion poten�tial shifts towards more positive potentials by 100–200 mV, but remains relatively negative. Therefore,the rubbing of surfaces does not promote corrosion,i.e. tribocorrosion of such alloys might be relativelyweak. When tungsten alloys are exposed in open air for6 months, the corrosion potentials shifts towards pos�itive values by 0.4–0.5 V compared with values of cor�rosion potential of freshly electrodeposited alloys. Thetribocorrosion testing was carried out on electrode�posited nanocrystalline Fe–W (26 at % W), Co–W(28 at % W); Co–P–W (4.5 at % P, 20 at % W) alloysexposed to ambient air for 6 months. Based on corro�sion potential measurements and electrochemicalimpedance data, the corrosion resistance under slidingconditions in 1M Na2SO4 at 5 N load decreases in thesequence: Fe–W > Co–W > Co–P–W [146].

The corrosion behavior of nanocrystalline (“amor�phous�like”) tungsten alloys is intriguing, as one canexpect that the increase of the amount of tungsten inthe coating will increase its corrosion resistance. Ni–W alloys electrodeposited from a citrate solution con�sist of three different phases, which are all present inhigh W�content alloys: a solid solution of W in a Nimatrix; an intermetallic compound Ni4W, as well asanother solid solution of W in Ni, with a W contenthigher than 20 mol %. The XPS analysis revealed thatthe alloys were covered with a surface layer of a com�plex structure containing pure metals Ni and W,Ni(OH)2 and WO3, on the very surface, as well as somepartially reduced oxide WO2.72 (most probably a solidsolution of WO2.72 in Ni) and tungsten carbide in the

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layer underneath. It is highly probable that some of theoxide species in the layer act as intermediates in thecathodic deposition process. Identifying these speciesshould be the clue to a more detailed understanding ofthe mechanism of induced deposition of W than hasbeen achieved so far. Corrosion of Ni–W alloys in sul�furic acid solution at corrosion potential occurs by thepreferential dissolution of nickel from the surfacelayer. The longer the corrosion process lasts, the morethe surface behaves like pure W. The lowest initial cor�rosion rates were recorded with alloys rich in W, butafter aging at OCP the lowest corrosion c.d. was foundfor the Ni–W alloy with the most homogenous phasestructure, consisting of the solid solution only [149].Based on the data obtained in neutral and acidic sul�phate solutions presented in [66, 110, 143, 148], it canbe stated that resistance to corrosion of nanocrystal�line tungsten�rich alloys decreases (i.e. corrosion cur�rents increases) in the following order: Ni–W; Co–W;Fe–W. However, there are still some ongoing contro�versial issues that deserve further investigation.Namely, while the experimental data presented in[145, 150] showed that the best corrosion resistancewas achieved in the cases of Co–W and Ni–W coat�ings with a lower content of tungsten, other authors[110] demonstrated that all tested Ni–W coatingscontaining 11–21 at % W exhibited a quite compara�ble corrosion resistance in NaCl solution. Probably,the rate of corrosion of tungsten alloys depends onboth composition and structure of the deposited alloy.Amorphous alloys, whatever way they are prepared,are metastable. In turn, nanocrystalline materials havea large fraction of mass concentrated at grain bound�aries; therefore they are also in a metastable state dueto an excess of free energy stored at these grain bound�aries that can be reduced in the course of a higher localcorrosion rate. On the other hand, the alloys havinglarger grains corrode at lower rates. Consequently, onprogressive heating, an amorphous/nanocrystallinealloy will first lower its energy by relaxation to a morestable state and at a higher temperature it will tend tocrystallize, thus resulting in an increase of the grainsize. Therefore, the rate of corrosion after annealingamorphous�like alloys decreases, e.g Fe–W [151],Fe–Ni–W (with 55 wt % of W) [128]. It is also worthmentioning that the corrosion resistance of amor�phous�like tungsten alloys with iron group metals canbe improved by introducing phosphorous as an alloy�ing element. The corrosion potential is moved towardsmore noble values by adding P to the deposit, and thepassive current density then drops remarkably with theaddition of W. Hence, amorphous Ni–W, Ni–W–P,and Fe–W alloy deposits show very high corrosionresistance in acid solutions [152].

Magnetic Properties

Because of the outstanding mechanical, magneticand corrosion�resistance properties, electrodeposited

tungsten�containing alloys have been proposed inrecent years for a widespread variety of magneticapplications, ranging from recording media [29, 105]to remotely�actuated micro�/nano�electromechani�cal systems (MEMS/NEMS), such as microactuators,micromotors, sensors, microgears or micromechani�cal magnetometers [153–155]. Amongst the investi�gated films are: binary alloys (e.g., Co–W [153–159],Ni–W [29, 105, 160], Fe–W [160]), ternary alloys(e.g., Co–Pt–W [161] or Ni–Fe–W [162]) or evenmore complex systems (e.g., Co–Pt–W–P [161, 163]or Co–Ni–Re–W–P [164]). Magnetic properties ofthese alloys (coercivity, HC, anisotropy easy axis direc�tion, saturation magnetization, MS, squareness ratioor remanent�to�saturation magnetization ratio,MR/MS, etc.) depend not only on the compositionbut also on the micro�structure of the obtained films.

Both hard and soft ferromagnetic properties havebeen reported in electrodeposited Co–W alloys (seeTable 3), depending on the exact electrodeposition con�ditions (in particular, the applied current density, j). Inmany Co�based alloys (including Co–W), the micro�structure of films consists of isolated ferromagneticCo�rich grains surrounded by non�magnetic orweakly�magnetic boundaries, where alloying elements(e.g, W, P, Pt, Mn) tend to concentrate [153]. How�ever, Co–W films with a more uniform microstructurehave been also prepared by electrodeposition usingaqueous baths as electrolytes, with and withoutorganic additives [154, 156–159]. Alloys with tungstenweight percentages of up to 30–40% have been readilyobtained [155, 159]. Remarkably, while the alloy com�position does not depend significantly on the appliedcurrent density (j), the microstructure of filmsbecomes progressively more refined as j increases,leading to pronounced variations in the coercivity andsquareness ratio values. The formation of hard mag�netic hcp Co3W is usually promoted when films areprepared at low current densities [157], whereas larger,values render amorphous Co–W films with concomi�tantly lower coercivity, both along in�plane and per�pendicular�to�plane applied field directions. Thecoercivity of Co3W films also markedly depends on thedegree of texture [157]. Typically, a (001) texturedevelops during growth of Co3W hcp films, favoringthe occurrence of hard ferromagnetic behavior andperpendicular�to�plane magnetic anisotropy [155].Conversely, fcc and amorphous Co–W are magneti�cally softer [154, 157, 158]. It is also possible to pre�pare Co3W/Co–W multilayers, where HC and MR/MS

can be precisely tuned by varying the layers thickness[154, 158]. The exchange coupling between bothphases causes the formation of kinked hysteresis loops,typical of spring�magnet behavior [165]. As expected,the saturation magnetization of Co–W alloys tends todecrease with the W percentage [157]. However, a fur�ther reduction in MS is observed when films becomeamorphous [158].

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In addition to Co–W, other binary W�containingelectrodeposited alloys (e.g., Ni–W or Fe–W) havebeen also electrodeposited. These films can exhibiteither hard [29, 105] or soft [160] ferromagnetic prop�erties, depending on the microstructure. Soft ferro�magnetic properties (i.e., low HC and MR/MS ratios)have been reported in composite coatings consisting ofW particles embedded in Ni or Fe films [160]. Hardferromagnetic behavior, with perpendicular�to�planecoercivity values ranging between 250 and 1500 Oe,has been obtained in Ni–W films (W weight percent�age in the range 6–18%) consisting of columnarNi⎯W fcc grains of about 5–25 nm in lateral size, with(111) texture, embedded in a Ni–W amorphousmatrix. Similar to Co–W alloys, the saturation mag�netization of Ni–W films decreases with the W con�tent, not only because of the inclusion of non�mag�netic W atoms (dilution law) but also due to the pro�nounced reduction in the crystallite size that occursfor high W percentages [29, 105].

Enhanced hard magnetic properties (i.e., perma�nent magnet�“like” behavior) have been reported inCo–Pt–W–(P) electrodeposited alloys, with coerciv�ity values as high as 2450 Oe and 2720 Oe along the in�plane and perpendicular�to�plane directions, respec�tively [161, 163]. Again, texture effects seem to play animportant role in the observed anisotropy direction.While (001) crystal orientation favors perpendicular�to�plane magnetic anisotropy, the occurrence of otherpreferred orientations promotes in�plane anisotropy

easy axis. Interestingly, post�synthesis chemical treat�ments of Co–Pt–W–(P) films in H2SO4 solution havebeen found to be beneficial for the increase of the HC

and MR/MS values of this type of films, although atwo�fold decrease in the saturation magnetization isobserved [164]. Contrary to Co–Pt–W–(P), Ni–Fe–W alloys exhibit soft ferromagnetic behavior, with thecoercivity lower than 10 Oe and the saturation magne�tization up to 1 T [162].

Finally, hard ferromagnetic behavior has been alsoencountered in Co�rich electrodeposited Co–Ni–Re–W–P alloys, with Co and W weight percentagesaround 70% and 7%, respectively [164]. The coerciv�ity value in these films, measured along the perpendic�ular�to�plane direction, ranges between 500 and2500 Oe and the MR/MS ratio between 0.15 and 0.3,whereas hard�axis loops with low coercivity are mea�sured along the in�plane direction (Table 3). The per�pendicular magnetic anisotropy can be ascribed to theprominent (002) hcp crystallographic texture.Remarkably, due to the shape anisotropy, the square�ness ratio of this material can be further increased, upto MR/MS = 0.45, if the continuous films are pat�terned to form arrays of micro�cylinders with a 2 : 1length : diameter aspect ratio [164]. Aside from theobvious use of these hard�magnetic alloys in high�density recording media, other applications can bealso envisaged, such as microactuators and magneticsensors [164].

Table 3. Summary of magnetic properties* of different W�containing alloys

Composition Microstructure MS (Tesla) MR/MS HC (Oe) Reference

Co–W (6–44 wt % W)

Amorphous for high current density 1–1.5 0.2–0.8 (||) 20 (||)30(⊥)

[154, 157, 158]

Co3W (001) hcp textured films (⟨D⟩ = 10–60 nm), low/intermediate current density

~1.4 0.2–0.5 (||)0.3–0.7(⊥)

200–600(||)50–900(⊥)

[156–159]

Co3W/Co–W hcp Co3W + fcc/amorphous Co–W (spring�magnet); multilayers

1.2–1.3 0.2–0.5(||) 3–250 (||) [154, 158]

Ni–W Nanocomposite films 0.3–0.6 ~0.05 (||) ~10 (||) [160]

Fe–W Nanocomposite films ~1.8 ~0.5 (||) ~10 (||) [160]

Ni–W Columnar fcc grains 0.4–0.6 0.1–0.6 (||)0.1–0.8 (⊥)

160–650 (||)250–1500(⊥)

[29, 105]

Ni–Fe–W Nanocrystalline fcc grains ((⟨D⟩ = 25–40 nm)

0.6–1 ~0.05 (||) ~10 (||) [162]

Co–Pt–W–(P) Columnar hcp grains; ⊥ anisotropy favored by the (001) texture

0.8–1 0.05–0.5 (||)0.1–0.7 (⊥)

60–2450 (||)650–2720 (⊥)

[161, 163]

Co–Ni–Re–W–P Columnar hcp grains with (002) tex�ture

~1.2 0.05–0.1 (||)0.15–0.3 (⊥)

~100 (||)500–2500 (⊥)

[164]

Note: *) HC, MR and MS denote the coercivity, remanent magnetization and saturation magnetization, respectively. The symbols || and⊥ indicate results from hysteresis loops recorded along the in�plane and perpendicular�to�plane directions, respectively. ⟨D⟩ refersto the average crystallite size, as evaluated from x�ray diffraction experiments.

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APPLICATIONS OF TUNGSTEN ALLOYS

Hydrogen Electrocatalysis

One of the most attractive applications of W alloysis the electrocatalytic production of hydrogen. Hydro�gen is considered to be the energy vector for the futureand its production, storage, and transporta tion arecurrently receiving a great deal of attention. This hasprompted a lively research of new materials able to ful�fill the demands of the sustainability in the energeticsector. In this context, Ni–W alloys have been pro�posed as electrocatalysts for hydrogen productionsince they offer lower overpotentials for the hydrogenevolution reaction (HER) than pure Ni. This has, inturn, stimulated the research on the electrodepositionof Ni–W alloys and the assessment of their HER prop�erties. In principle, crystalline materials are preferredover amorphous ones because the former are moreactive for the HER due to a larger number of electro�chemically active sites on crystalline surfaces. In anycase, studies in this field are still not numerous and itis envisaged that the interest on the electrodepositionof W alloys will experience a revival in the nearfuture. The HER on electrodeposited Ni⎯W alloyshas been studied both in alkaline and acidic media[28, 166–170], and even when the Ni–W is depositedin situ with the catalyzed hydrogen [169]. Navarro�Flores et al. [28] considered crystalline Ni3.4W with ahigh intrinsic electrocatalytic activity, explained thatby the modification of the electron density in d�orbit�als upon alloying nickel with tungsten. It also havebeen observed that an increase in crystallinity leads toan enhancement of the HER. In constrast, accordingto Wang et al. [170] amorphous alloys electrodepositedunder a super gravity field exhibit enhanced HERactivity in 10% NaOH solution compared to the activ�ity of those deposited under normal gravity conditions.Those authors showed that the quality of Ni–W layershad important implications on the HER properties.While the layers deposited under normal gravity con�ditions consisted of nodular grains and featuredmicrocracks, those deposited under super gravityfields were fine�grained and crack�free. As a result, thecurrent density of the HER at an overpo�tential of⎯150 mV vs. SHE increased from 3.3 mA cm–2 to82 mA cm–2. Such an increase was attributed to boththe increase in surface roughness and the intrinsicactivity.

Moreover, Navarro�Flores et al. [28] studied theHER in 0.5 M H2SO4 solution at 22°C by the Tafel lin�ear polarization measurements for pure Ni and Ni3.4W(23 at % W) electrodeposits. They found out that thepure Ni coating exhibited the classical Tafelian behav�ior within the explored overpotential window(between –0.5 and 0 V vs. SHE), indicating that theHER was a kinetically controlled Volmer�type reac�tion. Conversely, the HER kinetics for the Ni3.4Wcoating was controlled by the Heyrovsky desorption atlow overpotentials and by the Volmer reaction at

higher overpotentials. The Tafel value thus increasedfrom 40 mV dec–1 to 123 mV dec–1. Such a change inthe Tafel slope was mainly attributed to the increasedability of Ni3.4W to form hydrides compared to pureNi, as demonstrated by electrochemical impedancespectroscopy (EIS) experiments. The authors in [28]estimated that the amount of energy (i.e., the overpo�tential) required to maintain a fixed current density of1 mA cm–2 and thus a fixed hydrogen production ratewas –379 mV for Ni and –85 mV for Ni3.4W (i.e., 80%lower power input). Indeed, the Ni3.4W electrodepositdisplayed a higher intrinsic electrocatalytic activity,likely due to an increased electron density aroundNi�sites.

The electrocatalytic performance toward the HERof a series of Ni�based alloys prepared by electroless dep�osition was studied [171]. Ni–W–P and Ni–Mo–Palloy electrodes exhibit good electrocatalytic activitytoward HERs, with some Ni–Mo–P alloys approach�ing the performance of Pt foils. The electrocatalyticactivity of Ni�based alloys decreases with increasing Pcontent but increases with increasing content of W orMo. The experimental results can be interpreted interms of a pronounced synergy between Ni, which hasinternally paired d�electrons, and W or Mo, whichhave empty or half�filled d�orbitals and therefore, bindhydrogen atoms strongly. The absorption of hydrogenin these alloys is also found to have an important effecton their activity. Electrodes with higher activity arecapable to incorporate larger amounts of hydrogenduring cathodic charging.

It has been recently demonstrated that the HERactivity of Ni–W electrodeposits can be enhanced ifthe two metals are codeposited with suitable metaloxide particles. In [172], Ni–W/TiOx composites werepotentiostatically deposited onto pristine carbonfibers and the HER was studied in 0.5 M H2SO4 at⎯500 mV vs. SCE by the EIS. The Ni–W/TiOx�coatedcarbon fibers showed lower charge transfer resistance(i.e., a higher HER activity) compared to that of theNi–W�coated fibers.

The usefulness of Ni–W electrodeposits as a cath�ode in microbial electrolysis cells (MECs) has beenalso investigated by Hu et al. in [173]. They demon�strated the hydrogen production in a single�chambertubular MEC equipped with a Ni–W alloy electrodepos�ited onto carbon�fiber�weaved cloth that served as thecathode. At an applied voltage of 0.6 V the MEC exhib�ited a hydrogen production rate of 1.5 m3/day/m3.Though the hydrogen production was lower whencompared to the Pt cathode (2.3 m3/day/m3), thisstudy demonstrated a great potential of Ni–W alloysas cathode materials in MECs. Also, the replacementof platinum metals in catalysts by Ni–W alloys is dis�cussed based on the exchange currents of hydrogenevolution reaction [174]. The results of this studyallow concluding that nickel�tungsten alloys contain�

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ing 15–35% deposited on nickel are the most perspec�tive catalytic materials of non�platinum group.

Compared to Ni–W, the HER in the electrodepos�ited Co–W alloy films has been scarcely reported[175–177]. The exchange current density (ECD) ofhydrogen evolution as a function of the content ofCo–W alloys was reported in [145]. It was determinedthat for hydrogen evolution as a function of the con�tent of Co–W alloys has a minimum corresponding totransitive from polycrystalline to nanocrystallinestructure (20–22 at % of W); the values of the Tafelcoefficient for Co–W alloys are in the range 90–130 mV dec–1. Aravinda et al. [81] showed that theHER in alkaline citrate�containing solutions (pH 8.0)was enhanced by an increase in the concentration ofeither tungstate ions or dimethyl sulfoxide (DMSO)during the deposition of a Co3W film. According to theauthors, the adsorption of DMSO molecules onto theelectrode surface would favor the electrochemical dis�charge of water since these contain sulfur, which is asa well�known reducing agent.

Deposition into Recesses Microstructures

Novel microdevices can be realized by mergingelectrodeposition with substrated, pattered with UV,lithography, high energy X�ray or e�beam lithography.The fabrication of new MEMS, made up of mechani�cal, fluidic, and electrical elements of micrometerscale, has the potential to change our way of life infuture as dramatically as microelectronics has trans�formed it today. Its range of dimensions gives a possi�bility of mixing nano�, micro� and macro�structuresfabricated with extreme precision with the same pro�cesses. The versatility of high aspect ratio microfabri�cation processes permits the manufacture of trulythree�dimensional microstructures. In order to realizethese structures deep imprints are made into a poly�meric resist and are filled with metals using elec�trodeposition. Precisely machined high aspect ratiostructures are thus possible. The choice of elec�trodeposition as a method for depositing metals intomicrostructures is motivated by the low capital invest�ment requirements, and elimination of subtractivepatterning procedures. The resulting metallic struc�ture is used either as the final product or as a mold forplastic processing.

Single metals or alloys can be electrodeposited intodeep recesses for MEMS production. Alloy depositionis inherently more difficult to control than single metaldeposition but it offers a wider array of tailored mate�rial properties. A review of a variety of metal alloysdeposited into MEMS structures is given by Ehrfeld,et al. [26], including their proprietary structures withNi–W.

Recently, an increasing interest in electrochemicalalloy deposition has emerged as a result of novel micro�fabrication techniques for micromachines by the LIGA

process [178–180], (LIGA is a German acronym thatstands for: Lithography, Galvanoformung—Elec�trodeposition, Abformung—Plastic Molding), and gen�erally refers to the use of high energy x�rays, producedat a synchrotron, in resist exposure for the lithographicstep. The LIGA process was one of the first high aspectratio microfabrication processes, which was princi�pally developed at the Institute for MicrostructureTechnology at the Nuclear Research Center,Karlsruhe, Germany (IMT�FzK) in response to theneed for precisely machined nozzles having highaspect ratios, principal dimensions of tens ofmicrometers, and precise vertical sidewalls [178, 179].Subsequent research at the IMT�FzK, the Universityof Wisconsin and other laboratories has further refinedthe process and produced numerous device proto�types, e.g., see [181, 182]. Sharp corners and verticalsidewalls are obtained; deviations from the vertical ofonly 0.055 μm/100 μm over a height of 400 μm havebeen verified [183]. Resist layers up to ten centimetersthick have been patterned, making the fabrication ofstructures with aspect ratios much greater than 100 : 1possible. The exposed and developed resist can be sub�sequently used as a substrate for an electroplating pro�cess. The metal structure can also be used as a moldinsert, for plastic molding, to fabricate large volumesof plastic micro�scale parts. Tungsten alloys may be achoice material to fill this niche in developing micro�molds and multiscale molds.

Reports of tungsten alloy plating in LIGA micro�structures have generally centered on electrodeposi�tion of tungsten alloys using known electrolytes andmeasuring the resulting properties; rightfully so, as theproperty is the key parameter in adopting tungstenalloys. As an alternative to LIGA Ni, particularly forhigh�temperature MEMS applications, the tungstenalloy can greatly improve microstructural instabilitiesand loss of mechanical strength at intermediate tem�peratures when compared to Ni microstructures[183–188]. Replacing electrodeposited Ni with anelectrodeposited Ni–W alloy has improved MEMSstructures at elevated and cycled temperatures asnoted and reviewed by Suresh et al. [189] and Haj�Taieb et al. [190], but with lower ductility, fracturetoughness and increased internal stress than Ni. Inboth reports [189, 190] electrodeposited Ni–Wmicrotensile specimens with a thickness of 150 μmprepared using the LIGA process, and deposited froman ammonium citrate bath. Suresh et al. [189] showeda Hall�Petch plot of the measured yield strength ofNi–W samples as a function of the inverse square rootof grain size with a marked increase in yield strengthwith the addition of 5 and 15 at % W and found outthat the underlying micro�structures of the Ni–Walloys are stable up to 700°C. Those researchers con�cluded that as�deposited Ni–W exhibits brittle failure,but annealed LIGA Ni–W samples had a superior mixof thermal stability, strength and ductility. Haj�Taiebet al. [190] investigated various annealing tempera�

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MODERN TRENDS IN TUNGSTEN ALLOYS ELECTRODEPOSITION 511

tures. They observed a small grain growth, compared toelectrodeposited LIGA Ni deposits, with micro�hard�ness values of the LIGA Ni–W higher than those of thepure LIGA nickel. The 15 at % W structures exhibited asmaller decrease of the hardness with an increase of theannealing duration. Tensile tests reported the ultimatetensile strength (UTS) above 750 MPa, but lower duc�tility than pure nickel. Armstrong et al. [191] were thefirst to design a methodology to measure fracturetoughness of micro�scale specimens. They electrode�posited Ni�W alloys from an ammonia�citrate electro�lyte [52, 108] with beam structures having nominaldimensions of 60 μm × 20 μm × 14 μm. The loadapplied using a nanoindenter and fracture toughness,calculated from the fracture load was in the range of1.49–5.14 MPa√m for a Ni–W structure containing12.7 at % W. This is a higher fracture toughness of Si(another important MEMS material), but unfortu�nately considerably lower than that of electrodepos�ited nickel and other nickel�based alloys.

Electrodeposited Ni–W into recesses can be prob�lematic if there is high internal stress, and microstruc�tures that can potentially be tens to hundreds ofmicrons tall could develop cracks. Several investiga�tors made reports on using sulphamate nickel salts inammonium�citrate electrolytes: Wang et al., [192] andNamburi and Podlaha et al. [193], and from boricacid�citrate, non�ammonia electrolyes, Slavechevaet al. [27], for the purpose of reducing internal stress.It is interesting to note that the authors of these twopapers have observed a significant decrease of W % inthe deposit compared to thin films. This can indicatethe problematic nature of species under diffusion con�trol with an enhanced boundary layer due to the depthof the recess. In particular, even if the reacting metalion species are not under diffusion control, the result�ing hydroxyl ion produced by the deposition reactioncan be. In this case, the local pH can increase and alterthe deposit composition.

An example of a tall, 500 μm Ni–W deposit froman ammonia�citrate electrolyte is shown in Fig. 10.The electrolyte contained 0.25 M sodium citrate,0.4 M sodium tungstate, 0.2 M nickel sulfate and suf�ficient ammonium hydroxide to obtain a pH of 10.The deposit was pulse plated between –17.5 mA/cm2

and current density with long pulse times, with on�time = 15 s and off�time = 45 s. With this particularchoice of electrolyte and current density, dc deposi�tion at the same current density did not result in adeposit, but a clear precipitation of salts was visuallyobserved, indicative of a large surface pH rise. Key tothe design of the pulse plating conditions is that suffi�cient time is chosen to match the diffusion time scalesof the hydroxyl ion from the bottom of the recess.Hence, the pulse off�time is on the order of many sec�onds as opposed to fractions of a second for conven�tional pulse plating. The composition was measured at20 equidistant points along the length of a micropost,with the wave�length dispersive X�ray analysis (WDS),

and the composition was found to be uniform but low,having 3 wt % W along the direction of growth.

Nanowires of electrodeposited alloys can be readilyfabricated through templating techniques into nanop�orous membranes,which has been widely demon�strated, mostly in the area of magnetics [194–196]. Aparticular niche of the electrodeposition process isthat it is the preferred, and in some cases the only, pro�cess that renders the deposition of multilayerednanowires, which has been used to deposit giant mag�netoresistant materials (GMR), where a change in themagnetic field changes the deposit resistance. Nanow�ires with nanolayers of a ferromagnetic magneticmaterial separated by a non�ferromagnetic material,such as Cu, can easily present the current�perpendic�ular to the multilayer plane (CPP) type of GMR asopposed to the current�in�plane (CIP) GMR for thinfilms, providing more sensitivity [197–199]. There arenumerous electro�deposition examples of nanowire,multilayered alloys in literature, such as systems withCo–Cu/Cu [200–205], Co–Ni–Cu/Cu [206–208],as well as electrodeposited alloys nanowire arrays thatare unlayered for non�GMR magnetic applications,for example, Co–Ag [209, 210], Ni–Fe [211] Ni–Co[212], Ni–Cu [213] and Ni–Fe–Co–Cu [214].

Nanostructures

In a study to develop a Cu–Ni–W nanowire alloywith the ability for composition modulation, Guptaand Podlaha [215] used an ammonia�citrate electro�lyte with 60�m thick, commercially available anodizedaluminum oxide membranes (Whatman, Anodisc 25),with a listed pore diameter of 20 nm. Multilayeredalloys with one layer rich in Cu and the other layer richin Ni–W were fabricated. The electrolyte contained0.6 M Na3C6H5O7, 0.2 M Na2WO4, 0.3 M NiSO4 andvariable CuSO4 concentrations at the pH = 8, adjustedwith ammonium hydroxide at 70 ± 2°C. Since copperis more noble than nickel or tungsten, it can be prefer�entially deposited at a low current density; and an

100 μm

500 μm

Fig. 10. Example of a tall, 500 micron Ni–W structurewith 3 wt % W [193].

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alloy rich in Ni–W with some Cu can be deposited ata high current density. Due the presence of ammoniain the baths, the reversible potential of copper isshifted to a more negative potential range, hinderingthe ability to form multilayers with a pure Cu layer.Without ammonia, there is a larger separation of thereversible potential of copper and nickel�tungsten,that facilitates multilayer fabrication, however a lowerammonia concentration or no ammonia lowers thecurrent efficiency appreciably for nickel�tungstencodeposition without addition of an alternative buffer,such as boric acid. Moreover, the nanowire depositionwas confounded by the formation of oxide during themodulation, and results therein recommended thatthe potential of a more noble step be more negativethan –0.9 V vs SCE to avoid this situation.

Bairachna, et al. [216] electrodeposited Ni–Walloys from a non�ammonical electrolyte, containingcitrate and boric acid. Deposition was carried out inpolycarbonate membranes that are much thinner,hence provoking less of a change in transport bound�ary conditions when compared to thin films than thealumina used in the previously cited study. Nanowireswere successfully deposited with dc or pulsed currentdensity. A typical example of released nanowires usingdichloromethane to dissolve the membrane is shownin Fig. 10, with two types of measurements. The diam�eters of the nanowires are between 135–155 nm due tonondiscrete edges. The surface of the nanowiresappears to be modulated, that is irrespective ofwhether the deposits are dc or pulse plated, and someof the measurements in Fig. 11 show the size of bumps(~30–40 nm). The modulation effect may be similar

to the laminations first noted in thin films by Brenneret al. [6] in 1947, that also occurred when the currentwas not pulsed or interrupted.

Cesiulis et al. [67] developed conditions to elec�trodeposit Co–W–P and Co�W into nanoporous alu�mina oxide membranes, utilizing an electrolyte thatcontained nickel sulfate 0.15 M, sodium tungstate0.10 M, trisodium citrate 0.50 M, boric acid 1.00 M,sodium hydroxide or sulfuric acid to adjust pH to 7.00,as well as with and without 0.05M NaH2PO2. Theelectrodeposition was performed under constantpotential and pulse potential deposition modes, andthe current was monitored to define the end of the wireor tube growth. The amplitude of the potential in bothmodes was chosen based on thin film deposit resultscorresponding to a potential value that should beattributed to the current density of ~15 mA cm–2 atsteady state (the approximation is due to an estimateof membrane porosity.) At this current density, themaximal content of W in the alloys is obtained and thecurrent efficiency is the highest, reducing the gaseoushydrogen evolution. In pulse potential deposition, thevalue for a relaxation potential (current equal to zero)was set as –0.66 V. The pulse duration was 10 s, therelaxation duration was 30 s, and the total number ofcycles was 1500. Interesting results showed that whilenanowires could be obtained with pulsed deposition,nanotubes were observed in a dc mode, Fig. 12. Nan�otubes have also been observed in electrodepositedCo�Ni�Fe�Cu nanotemplated deposits [217, 218]and, similarly, Bi–Te [219]. The tube formation wasattributed to the directed potential field along themembrane wall due to mass transport effects, as a con�sequence of hydrogen evolution. The influence of theside reaction and electrolyte pH was also cited byFukanaka [220] in the electrodeposition of Ni nano�tubes. Nanotubes were observed by Verbeeck [221],too, when depositing Fe nanowires. If short depositiontimes were used, tubes were formed. They also ana�lyzed the structure and found oxide skins on the nano�

100 nm

24.4 nm

22.6 nm

138 nm

37.7 nm

132 nm

38.7 nm

140 nm

35.9 nm

34.3 nm

38.5 nm38.0 nm

39.7 nm

35.9 nm

157 nm

35.4 nm

153 nm

38.9 nm

153 nm

22.8 nm

24.8 nm

22.2 nm

125 nm

21.1 nm

22.5 nm

23.8 nm

21.9 nm

141 nm

Fig. 11. TEM of pulsed current Ni–W deposit from What�man polycarbonate membranes with a stated pore diame�ters of 50 nm at one end; actual pore diameter variesslighlty along length, ω (W) = 50.4 ± 0.5 wt %. (with per�mission from the Electrochemical Society) [206].

2 μm

Fig. 12. SEM of Co–W electrodeposits obtained from acitrate�boric acid electrolyte at a constant potential, E =–1.1 V vs SCE [67].

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MODERN TRENDS IN TUNGSTEN ALLOYS ELECTRODEPOSITION 513

structures. The oxygen source was thought to comefrom the polycarbonate membrane [220, 221].

ACKNOWLEDGMENTS

This work was supported by the FP7 projectsNANOALLOY (No. 252407) and TEMADEP(No. 05�104�7540), the MAT2011�27380�C02�01project from the MINECO (Spain), the US NationalScience Foundation under Grant No. 0957448, andthe Lithuanian ESF agency grant (code VP1�3.1�SMM�08�K�01�014).

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