Surface & Coatings Technology 307 (2016) 978–1010

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Surface & Coatings Technology

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Ni-W electrodeposited coatings: Characterization, propertiesand applications

M.H. Allahyarzadeh, M. Aliofkhazraei ⁎, A.R. Rezvanian, V. Torabinejad, A.R. Sabour RouhaghdamDepartment of Materials Engineering, Faculty of Engineering, Tarbiat Modares University, Tehran, P.O. Box: 14115-143, Iran

⁎ Corresponding author.E-mail addresses: maliofkh@gmail.com, khazraei@mo

http://dx.doi.org/10.1016/j.surfcoat.2016.09.0520257-8972/© 2016 Elsevier B.V. All rights reserved.

a b s t r a c t

a r t i c l e i n f o

Article history:Received 13 June 2016Revised 20 September 2016Accepted in revised form 23 September 2016Available online 27 September 2016

Nowadays, application of nanocrystalline nickel-tungsten (Ni-W) alloys is receiving a great interest because theyare an efficient replacement for hard chromium coatings owing to their premium hardness, wear, and corrosionproperties. Moreover, heat-treated nanocrystalline Ni-W alloys demonstrate proper mechanical properties aswell as thermal stability at high temperatures. The current paper reviews different aspects of electrodeposition,microstructure, corrosion, oxidation, wear, and mechanical properties Ni-W alloys and nano/micro composites.Besides, heat treatment effects on properties and thermal stability of these alloys are also reviewed.

© 2016 Elsevier B.V. All rights reserved.

Keywords:Nickel-tungstenElectrodepositionMicrostructureMechanical propertiesHeat treatment effectsCorrosionWear

Contents

1. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9791.1. Why nickel-tungsten coatings? . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9791.2. Application and properties of nickel-tungsten coatings . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 979

2. Different types of utilized electrolytes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9802.1. Electrolytes based on Watts bath . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9802.2. Sulfate-citrate containing electrolytes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9802.3. Sulfamate-citrate electrolytes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9802.4. Other electrolytes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 981

3. Mechanism of Ni-W electrodeposition . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9814. Effect of different plating parameters . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 982

4.1. Electrolyte chemical composition . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9824.2. pH effect . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9834.3. Temperature . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9834.4. Current density . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 984

4.4.1. Micro/nanocomposite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9844.5. Rotation rate . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9844.6. Different additives . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 984

5. Techniques for Ni-W electrodeposition . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9855.1. Direct current technique . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9855.2. Pulse current technique . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 985

5.2.1. Pulse duty cycle . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9865.2.2. Pulse current density . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9865.2.3. Pulse frequency . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 986

5.3. Pulse reverse technique . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9866. Microstructural properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 987

dares.ac.ir (M. Aliofkhazraei).

979M.H. Allahyarzadeh et al. / Surface & Coatings Technology 307 (2016) 978–1010

7. Mechanical properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9887.1. Microhardness and Hall-Petch equation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9887.2. Tensile strength . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 989

7.2.1. Residual stresses in Ni-W electrodeposits . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9907.3. Creep and fatigue . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9917.4. Fracture toughness . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 992

8. Magnetic properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9929. Oxidation and high-temperature corrosion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 99410. Corrosion properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 995

10.1. Micro/nanocomposite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 99511. Tribological and wear properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 996

11.1. Micro/nanocomposite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 99712. Heat treatment . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 999

12.1. Microstructural evolution . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 99912.2. Heat treatment effects on mechanical properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 100012.3. Thermal stability . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1002

13. Modern trends in fabrication of novel Ni-W films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 100313.1. Diffusion barrier . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 100313.2. Ni-W alloys and nanocomposites with gradient structures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 100313.3. Ni-W alloys and nanocomposites with multilayer structures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 100313.4. Hydrophobic coatings . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1004

14. Mathematical modeling and computer simulation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 100415. Conclusions and future research directions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1005References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1006

1. Introduction

1.1. Why nickel-tungsten coatings?

Chromium-plated coatings have found a wide range of engineeringapplications because of their excellent corrosion and wear resistanceproperties [1,2]. In addition to their high resistance to corrosion andwear, chromium-plated coatings possess other useful characteristicssuch as high hardness, low friction coefficient, and appearance. Therealso exist some undesired facts about these coatings like low cathodiccurrent efficiency and low throwing and covering power which causehigh energy requirement for electroplating [1]. In spite of many benefi-cial properties and low production costs, the electrolyte used for platingcontains toxic hexavalent chromium ions that must be relinquished ac-cording to EU directives (2000/53/WE and 2011/37/UE); therefore, thisannouncement provoked researchers to put much effort into finding analternative for chromium [3,4]. Besides, it is obvious that hard chromi-um plating leads to degrading the fatigue resistance of engineeringparts due to residual tensile stresses and cracks. In this regard, nickel-based coatings alloyed with other refractory metals including tungstenandmolybdenum are introduced as promising and appropriate alterna-tives for their distinguished mechanical and thermal properties (highhardness and wear, corrosion and heat resistance). For this purpose,some researchers have fabricated nanocrystalline nickel-tungstenalloys in order to study their properties and characteristics. Ni-Welectrodeposition not only preserves the environment and avoidshazardous consequences, but also the produced coatings have highercorrosion and wear resistance compared to electroplated hardchromium coatings [5,6]. It is worth mentioning that traditionalalloying processes cannot be employed in the production of nickel-tungsten alloys due to the enormous difference in their meltingpoint (nickel: 1445 °C and tungsten: 3410 °C) as well as low solubil-ity. However, electrodeposition of these alloys by means of certainelectrolytes is an economic and simple method which enables the re-searchers to apply a uniform and proper nickel-tungsten coating ondifferent parts. It is also important to note that the released hydrogenduring nickel-tungsten electrodeposition could diffuse through thesubstrate lattice and cause failure of loaded specimens due to hydro-gen embrittlement [7].

1.2. Application and properties of nickel-tungsten coatings

Because of the high hardness andwear resistance of nickel-tungstenalloys, they are mainly used as hard and wear-resistant coatings. In ad-dition, the proper corrosion resistance of these alloys hasmade their ap-plication possible as hard coatings with noticeable resistance againstcorrosion [3,8–21]. Besides, it was reported that Ni-W films possess de-sirable electrocatalytic properties [22–26]. Hence, Ni-W, Ni-Mo, and Ni-W-Mo alloy systems are known as proper catalysts for hydro-sulfuration and hydrogenation processes of various organic compoundsandpetroleumproducts [23]. For instance, Hristova et al. [23,27] appliedNi-Mo-W alloys in the oxidation of sulfides to sulfur in alkaline solu-tions. In otherworks, Ni-W alloyswere used as catalysts in hydrogen re-duction processes and alkaline water electrolysis with successfulelectrocatalytic effects [22,24–26,28–34]. Tasic et al. [25] reported thatNi-Wcoatings obtained in baths of higher pH values demonstrate betterelectrocatalytic properties upon which hydrogen evolution potential isrelatively low. Besides, it was expressed that deposition current density,tungsten content, and other factors affecting the morphology of coat-ings are important in the enhancement of catalytic efficiency. For exam-ple, it is reported that compact coatings show better catalytic propertiesfor hydrogen evolution reaction in comparison with those with island-formed morphologies [29].

With rapid developments in the field of micromechanical technolo-gies, demands for excellent physical and mechanical properties contin-ue to grow; therefore, nickel-tungsten alloys have become an idealchoice in the manufacture of microelectromechanical structures andmicrodevices for their combined desirable hardness and wear resis-tance properties [35–39]. In this regard, the thermal stability of nickel-tungsten coatingswas studied byHaj-Taib et al. [35] and itwas reportedthat after annealing at 700 °C, the film demonstrates a stable structurewith a slight grain growth. These alloys were applied on aluminumfoams and resulted in an improvement in strength and a higher energyabsorption [40]. In electrical contacts, it is necessary to have an effectivebarrier layer against corrosion and diffusion. Because nanocrystallineNi-W can improve the corrosion resistance, minimize contact wear,and significantly reduce the formation of intermetallics resulting frominter-diffusion at interfaces, this alloy (Ni–15.8 at%W)may be nominat-ed as a reasonably effective barrier in electronic/electrical contacts [41].

980 M.H. Allahyarzadeh et al. / Surface & Coatings Technology 307 (2016) 978–1010

Although wettability features of coatings strongly depend on morphol-ogy and surface chemistry of coatings; electrodeposited nickel-tungstenalloys are also reported to be hydrophilic [42]. In particular, Ni-W-B al-loys have been successfully employed for reduction of coal processes[43]. Various applications of the nickel-tungsten alloy are schematicallyshown in Fig. 1.

2. Different types of utilized electrolytes

2.1. Electrolytes based on Watts bath

Several researchers have employed the well-known Watts bath toproduce nickel-tungsten coatings. As mentioned in the literature,Watts's bath commonly contains 300 g lit−1 nickel sulfate, 60 g l−1

nickel chloride, and 40 g l−1 boric acid along with some additivessuch as brighteners and stress relievers. Deposition conditions inWatts's bath are as: cathodic current density of 3–5 A dm−2, pH of 4–5.5, and temperature of 45–60 °C. A patent was registered for the syn-thesis of nickel-tungsten plated coatings in which sodium tungstate isintroduced toWatts's bath [44]. It was reported that addition of sodiumtungstatewould lead to precipitation of nickel- and tungsten-consistingsediments. These sediments are dissolved in the electrolyte by the addi-tion of sulfuric acid and reduce pH to about 3. It was reported that in so-lutions containing tungstate and nickel chloride or sulfate (without any

Fig. 1. Various applications o

ligand), the solubility and precipitation of hydroxides and tungstatesstrongly depends on the pH values [45,46]. Moreover, apart fromsodium tungstate, the addition of sodium citrate was reported in somestudies [11].

2.2. Sulfate-citrate containing electrolytes

Citrate baths aremainly used in the study of Ni-W electrodeposition[47–50]. In these electrolytes, nickel sulfate is commonly used as sup-plying resource of nickel ions, sodium tungstate is used to producetungstate ions (tungsten), and trisodium citrate is applied as acomplexing agent for Ni and W ions. Since ammonium chloride isemployed to increase cathodic current efficiency, these types of electro-lytes are named as citrate-ammonia electrolytes. Sodium bromide(NaBr) could often be added to increase bath conductivity and also re-duce crack formation [51]. Diluted sulfuric acid and ammonia solutionacid may be used in order to adjust the pH of these electrolytes.

2.3. Sulfamate-citrate electrolytes

Use of sulfamate-citrate electrolytes was also reported in Ni-W elec-trodeposition. Bratoeva and Atanasov investigated pH changes in therange of 4–8 in a solution bearing 16.5 g l−1 nickel sulfamate, 30 g l−1

sodium tungstate and 90 g l−1 trisodium citrate at a temperature of

f nickel-tungsten alloy.

981M.H. Allahyarzadeh et al. / Surface & Coatings Technology 307 (2016) 978–1010

60 °C and a current density of 5 A dm−2. They reported that the solutionpH significantly affects coatings characteristics such as tungsten con-tent, microhardness, and crystallite size. The pH of 7 was suggested asthe optimum value in ref. [52]. By applying sulfamate baths, residualtensile stress in nickel-tungsten coatingsmay be highly reduced.Matsuiet al. [53,54] employed a sulfamate bath containing nickel sulfamate,nickel chloride, sodium tungstate, and propionic acid (as complexingagent) at the pH value of 4.

2.4. Other electrolytes

Gluconate electrolyte is also used in electroplating of Ni-W films.This bath consists of 0.11 M nickel sulfate, 0.05 M sodium tungstate,0.5 M sodium gluconate, and 0.65 M boric acid. The pH is adjusted at 5and coatings are applied at 80 °C [55]. A bath with glycine as acomplexing agent for nickel-tungsten electrodeposition was used byMizushima et al. [56–58]. They introduced another bath withtriethanolamine as a complexing agent. In the same work, a bath con-taining triethanolamine and glycine mixture and a bath with threecomplexing agents (citrate + triethanolamine + glycine) were alsostudied. They stated that the most cathodic current efficiency and thehighest tungsten content were achieved in the bath with a mixture ofall three complexing agents. Recently, citrate-glycinate baths wereemployed in electrodeposition of Ni-W-B coatings [59]. Electrolytescontaining citric acid and glycolic acid (as complexing agent) were in-troduced and it was reported that coating obtained in these baths pos-sess noticeably higher wear and corrosion resistance and hardnessthan those produced in citric acid baths. It was also mentioned thatpresence of glycolic acid might cause improvement in cathodic currentefficiency [60].

Table 1Brief summary of Ni-W deposition conditions, deposition rates, efficiencies, and other data rela

Composition Current type Electrolyte Current d

1 Ni-(8–13)%atW DC Gluconate 4.7 A/dm(iave)PC Gluconate

2 Ni-20%at W DC Citrate-ammonia 5 A/dm2

Ni-25%at W 20 A/dm2

3 Ni-W DC Citrate- glycine 1–5 A/dm4 Ni-0.72% at W DC citrate-ammonia-boric acid 5 A/dm2

Ni-9.33% at W 20 A/dm2

5 Ni-W PRC Citrate-ammonia 20 A/dm2

~10.9 A/d6 Ni-7.4%at W DC Citrate-ammonia 10 A/dm2

Ni-17.5%at W7 Ni-13.3 at%W DC Citrate-ammonia 1.5 A/dm

Ni-23.5 at%WNi-33.5 at%W

8 Ni-(24–36)%atW DC Citrate 1 A/dm2

Ni-(24–36)%atW 1 A/dm2

Ni-(24–36)%atW 1.5 A/dmNi-(24–36)%atW 1.5 A/dm

9 Ni-12.4%atW DC Citrate-ammonia 15 A/dm2

10 Ni-9.3%atW DC Citrate-ammonia 20 A/dm2

Ni-16.1%atW 20 A/dm2

Ni-0.7%atW 5 A/dm2

Ni-2.4%atW 5 A/dm2

11 Ni-W DC Citrate 4 A/dm2

12 Ni-W DC Citrate-boric acid 1 A/dm2

5 A/dm2

Ni-Mo-W 1 A/dm2

5 A/dm2

13 Ni–W/ZrO2 DC Citrate-ammonia 5 A/dm2

11 A/dm2

14 Ni-W/diamond DC Citrate-ammonia 10 A/dm2

15 Ni-W-PCTFE (4 g/L) DC Tartrate –Ni-W-PCTFE (8 g/L)Ni-W-PCTFE (20 g/L)

Table 1 briefly represents the deposition conditions, cathodic cur-rent efficiencies, deposition rates, and other data related to some Ni-W electrodeposition.

3. Mechanism of Ni-W electrodeposition

According to Brenner classifications, nickel-tungsten alloys plating isconducted in the formof “induced co-deposition” [73]. Tungsten andmo-lybdenum ions could not be able to fully reduce in their aqueous solu-tions while they may be utterly reduced in the presence of iron groupelements and electrodeposit in the form of alloy on the cathode surface.Indeed, reduction behavior of somemetals such as tungsten andmolyb-denum in the presence of iron group elements including iron, nickel,and cobalt is called induced codeposition. Holt and Vaalar developedthe first model for induced codeposition [74]. According to that model,the following reactions occur on the cathode surface. In order to prog-ress in reaction II, it is primarily necessary that a thin film of productsof reaction (1) form on the surface of cathode acting as catalysts for re-action (2). Besides, with progress in reaction II and covering of cathodeby tungsten, reaction (2) is ceased. Further continuation of reaction (1)leads to creation of new sites catalyzing reaction (2):

Mþþaqð Þ þ 2e− →M sð Þ ð1Þ

WO−4 aqð Þ þ 8Hþ þ 6e−→

Mð ÞW sð Þ þ 4H2O lð Þ ð2Þ

In many studies, obtained XRD spectra demonstrate a solid solutionstructure of tungsten in nickel for as-deposited nickel-tungsten alloyswhich may be a great challenge for Holt and Valaar's model. Clark and

ted.

ensity pH Temp. (°C) Deposition rate (μm/h) Efficiency Ref.

2 5.0 80 ~35 μm/h ~61% [55]5.0 80 ~28 μm/h ~47.5%8.5–9.2 80–85 19.2 mg/cm2 h – [61]

60.3 mg/cm2 h2 7.5 60 – 60–85% [57]

8.5 75 ~23 μm/h ~35% [62]~85 μm/h ~31%

(iave) 7.4 75 40 μm/h – [63]m2(iave) 23 μm/h

6.5 50 – ~60% [48]~40%

2 8.5 – 31.2 μm/h 40% [64]24.9 μm/h 25%18.5 μm/h 17%

8 50 °C – ~15% [65]70 °C ~20%

2 50 °C ~12.5%2 70 °C ~18%

– 72 °C ~1 μm/min – [66]8.5 75 °C ~82 μm/h ~31.3% [67]

85 °C ~95 μm/h ~31.7%75 °C ~20 μm/h ~35%85 °C ~20 μm/h ~36.8%

4 20 °C – 24.4% [68]40 °C 28.2%60 °C 29.9%

7 25 °C – ~12% [69]~5.5%~7%~2.6%

8 60 °C – 18% [70]37%

8.5 75 °C 16.5–19 μm/h – [71]5 45 °C ~20 mg/hcm2 – [72]

~15 mg/hcm2

~10 mg/hcm2

982 M.H. Allahyarzadeh et al. / Surface & Coatings Technology 307 (2016) 978–1010

Lietzke [75] stated another procedure for electrodeposition of tungstenalloys on cathode: First, a partially reduced tungstate film is formed onthe cathode (kind of tungsten oxide) and then, in the presence of depos-ited nickel on the cathode, the tungstate film would undergo a catalyticcomplete reduction by hydrogen.Without the presence of inducing irongroup elements, tungstate ions could only be partially reduced and theirfurther and complete reduction would not be possible due to its lowoverpotentials to hydrogen evolution. Fukushima et al. [76] suggesteda mechanism for codeposition of Ni-Mo alloys which is very similar toelectrodeposition of Ni-W alloys. They stated that tungstate or molyb-date in form of oxide or hydroxide is initially reduced to W(IV)/Mo(IV) and then, with the adsorbed hydrogen, it would producemetal-lic tungsten or molybdenum in the presence of codeposited nickel [48].

According to the proposed model by Landholt and Podlaha, electro-chemical deposition of molybdenum or tungsten is conducted by forma-tion of an intermediate component that is adsorbed on cathode [77,78]:

NiCit− þWO2−4 þ 2H2Oþ 2e− → NiCitWO2½ �−ads þ 4OH− ð3Þ

NiCitWO2½ �−ads þ 2H2Oþ 2e− →Ws þ NiCit− þ 4OH− ð4Þ

NiCit− þ 2e− →Nis þ Cit2− ð5Þ

Based on the above mechanism, tungsten content in the alloy coat-ing may be increased by an increase in the concentration of tungstateand nickel intermediate compound in the boundary layer. Niu et al.[79] and Zeng et al. [80] showed the adsorbed nickel and tungstate (mo-lybdate) intermediate compounds by means of Raman spectrometry.

Younes and Gileadi investigated nickel-tungsten electrodeposi-tion with the following hypotheses in an ammonia-free electrolyte:1) Tungstate/citrate complex forms a ternary complex with nickelin solution or on cathode surface; 2) Nickel may deposit from itscomplexes with either citrate or with NH3 and 3) Tungsten (withnickel) may only deposit from a ternary complex with a general for-mula of:

Nið ÞP WO4ð Þq Citð Þmh i2 p−q−1:5mð Þ

While all stoichiometric coefficients were assumed to be 1, formulaof the ternary complex becomes simple:

Nið Þ WO4ð Þ Citð Þ½ �3−

The predominant complexes made of tungstate and citrate in pH =8 contains one proton and has the following formula:

Hþ þWO2−4 þ Cit3− → HWO4ð Þ Citð Þ½ �4− ð6Þ

Therefore, the formula of aforementioned ternary complex is:

Nið Þ HWO4ð Þ Citð Þ½ �2−

The ternary complex could explain the observation that althoughtungsten only codeposit with nickel, even if tungstate ions concentrationin the electrolyte are higher thannickel ions, a parallel route for nickel de-position from its own complexes (with citrate or NH3) always exist thatleads to increasing in nickel content in alloys. Therefore, removal of am-monia may contribute to increasing tungsten content in alloys [81,82].According to performed analyses, the concentration of this ternary com-plex was in the range of 2–4 mM. Due to its low concentration, the reac-tion of reduction of the protonated tungsten/citrate complexes iscontrolled by diffusion step. In order to complete reduction process,free nickel surface is also required [49,83]. However, Kabi et al. [84]showed that adsorption kinetic mechanism of the ternary complex maydepend on the deposition potential. Obradovic et al. [83] stated about

nickel-tungsten electrodeposition mechanism that nickel initially depos-ited from the ammonia-citrate complex. The ternary complex reductionwould be under activation control at lower overpotentials; while diffu-sion becomes the controlling step atmore cathodic potentials. Their stud-ies proved the existence of different complexes in the electrolyte; themost important of them which are necessary for Ni-W alloy depositionare protonated tungsten/citrate complex and nickel ammonia/citratecomplex. Spectrophotometry investigations revealed no complex bearingboth tungsten and nickel. At lower cathodic potentials, the presence oftungsten hydroxide/oxide layers control all the reactions on the cathodesurface. Sudden beginning of tungsten codeposition is attributed to acloser approach of tungstate/citrate complex ion from the freshly depos-ited Ni particles and electron tunneling from nickel particles attached tothe complex and reduction of tungsten to a metallic state.

Eliaz and Gileadi [85,86] reported that the adsorbed intermediate donot cause electrodeposition, but soluble complex species consistingnickel and tungstate ions [(Ni)(HWO4)(Cit)]2− are responsible for in-duced codeposition of Ni-W. The proposed mechanism by Podlahaand Landholt was modified in another article by introducing catalyticcomponents as responsible species for nickel-tungsten inducedcodeposition instead of adsorbed Ni(I) intermediate. This model is con-sistent with experimental data [69].

The basis of mechanism proposed by Sassi et al. [87,88] is themixedcomplex produced through the following reaction that leads to electro-deposition of nickel-tungsten alloy by taking eight electrons:

Ni C6H5O7ð Þ½ �− þ HWO4ð Þ C6H5O7ð Þ½ �4− ¼ Ni HWO4ð Þ C6H5O7ð Þ½ �2− þ C6H5O7½ �3−

ð7Þ

Ni HWO4ð Þ C6H5O7ð Þ½ �2− þ 8e− þ 3H2O ¼ NiþWþ C6H5O7½ �3− þ 7OH−

ð8Þ

Generally, via the mechanisms proposed by Fukushima et al.,Chassaing et al., Podlaha et al., and Sun et al. intermediate componentsin the form of adsorbed species which may contain metallic oxides re-sult in electrodeposition of Ni-W. However, soluble complexes in theelectrolyte are addressed as themajor responsible for electrodepositionof nickel-tungsten alloys in themechanismsproposed by Eliaz et al., Sunet al., Younes et al., and Sassi et al. Fig. 2 illustrates a brief schematic ex-planation for various proposed mechanisms.

4. Effect of different plating parameters

4.1. Electrolyte chemical composition

Categorizing electrodeposition of nickel-tungsten alloys as “inducedcodeposition” [73], electrolyte chemical composition may have an im-portant effect on the kinetics of Ni-W electroplating. According to theaforementioned mechanisms, tungsten may deposit from the interme-diate compound formedwith nickel at the cathode/electrolyte interface.It was indicated that there was more tungsten content in coatings ob-tained in electrolytes with higher concentrations of tungstate ionsthan nickel (tungsten-rich baths). Various studies demonstrated that ef-fects of deposition parameters such as pH, temperature, and currentdensity, in turn, depend on the bath chemical composition. Calculationsand experimental results of Cesilius et al. [89] illustrated that it wouldbe possible to achieve the highest cathodic current efficiency and depo-sition rate by altering the bath chemical composition and choosing op-timum concentrations for solution species. Wu et al. [90] investigatedthe effect of boric acid concentration in electrolyte (40 g l−1 H3Cit, 16g l−1 NiSO4, 32 g l−1 Na2WO4 and 50 g l−1 (NH4)2SO4) on Ni-W coat-ings electrodeposition and structure. They showed that by an increasein boric acid concentration from 0 to 1.5 M, tungsten content could beraised in the coating. However, cathodic current efficiency slightly in-creased in the beginning from 40 to 48% and slowly reduced thereafter

Fig. 2. Brief schematic explanation for various proposed mechanisms of Ni-W electrodeposition.

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(Fig. 3). It was observed that the effect of current density variations ontungsten content in the coating is affected by boric acid concentrationin the electrolyte. The increase and decrease in concentrations of Ni+2

and Cit3−, respectively, and the addition of nickel sulfamatewould facil-itate the increase in cathodic current efficiency [65]. In some Ni-Welectroplating bath, ammonia is used in order to adjust bath pH. Sinceammonia can act as a ligand for Ni complexes, it is possible to electrode-posit Ni-W films with High W content (50 at%) by removing ammoniafrom the electrolyte [81,82].

4.2. pH effect

A literature review of nickel-tungsten electroplating reveals that ef-fect of pH variation on tungsten content mainly depends on the electro-lyte chemical composition and deposition conditions. Variations in pHat different studies with different bath chemical compositions and depo-sition conditions offered distinct results. It was reported, in this regard,

Fig. 3. Boric acid concentration effect on the tungsten content and cathodic currentefficiency (at pH of 6.5 and temperature of 65 °C and solution of 40 g l−1 H3Cit, 50 g l−1

(NH4)2SO4, 16 g l−1 NiSO4, and 32 g l−1 Na2WO4) [90].

that pH had a distinctive influence on tungsten content, microhardness,crystallite size and cathodic current efficiency of nickel-tungsten coatingsobtained from sulfamate-citrate baths [52]. In fact, by increasing pHvalues from 4 to 7, tungsten content and microhardness are improved.In comparison, a further increase in pH value up to 8 led to a decreasein tungsten content and deterioration in microhardness of coatings. Themaximum amount of tungsten and microhardness was observed at thepH value of 7. Moreover, it was revealed that by increasing in pH valuefrom 4 to 6, cathodic current efficiency is also improved; however, a fur-ther increase in pH value from 6 to 8 caused dropping of cathodic currentefficiency. Younes et al. [91] reported different findings in W-rich citrateelectrolyteswhich are schematically demonstrated in Fig. 4. In ammonia-citrate bathswith equal concentrations of tungstate andnickel ions, tung-sten content and cathodic current efficiencywas improved by an increasein pH value from 7 to 9; yet, tungsten content and cathodic current effi-ciency were both diminished by a further increase in pH value up to 10[92]. Also, the increase in pH led to higher cathodic current efficienciesin electrodeposition of Ni-W-B coatings [93]. Studying the effect of pHvariations in baths containing boric acid and citrate Wu et al. [90] statedthat by increasing pH value from 5.5 to 7.5, cathodic current efficiencywas seen to rise; meanwhile, tungsten content was also slightlyincreased.

4.3. Temperature

There have been different, partly opposite, reports published aroundtemperature effects on nickel-tungsten coatings properties. Effect oftemperature on cathodic current efficiency, tungsten content, crystallitesize, and grain orientation was studied. Although it was reported thathigher temperatures normally result in higher tungsten content, finergrains and more random orientation in the Ni-W coatings [94–96],Younes et al. [65,81,82,97] stated that temperature had no significantimpact on tungsten content. Furthermore, most studies show that theincrease in temperature leads to a consequent increase in tungsten con-tent, faradaic cathodic current efficiency, and partial current of nickeland tungsten. It was reported that deposition temperature might affectthe structure and mechanical properties of nickel-tungsten coatings. Astudy carried out by Someskowa et al. [98] revealed that Ni-W coatings

Fig. 4. Effect of pH variation on tungsten content in Ni-W electrodeposited from electrolyte (a) containing 0.1 M nickel sulfate, 0.4 M sodium tungstate, and 0.6 M trisodium citrate andelectrolyte (b) comprising 0.05 M nickel sulfate, 0.4 M sodium tungstate, and 0.6 M trisodium citrate [91].

984 M.H. Allahyarzadeh et al. / Surface & Coatings Technology 307 (2016) 978–1010

produced at 80 °C were amorphous while those obtained at 75 °C dem-onstrated an amorphous-nanocrystalline structure. In addition, it wasstated that Young's modulus and hardness of coatings produced at75 °C were higher than those prepared at 80 °C.

4.4. Current density

Deposition current density is known to be influential in coatingschemical composition and structure; even though, the effect of currentvariations on nickel-tungsten coatings features is, in turn, affected bythe chemical composition of the deposition electrolyte [90]. Hou et al.[99] reported that in W-rich electrolytes, increasing the pulse currentdensity led to a higher tungsten content and microhardness in nickel-tungsten coatings. Besides, current density affects nickel-tungsten filmtexture [96] and seems to be substantially effective on surface morpholo-gy of nickel-tungsten coatings. It is reported that lower current densitieslead to the formation of smooth coatings with compact morphology;while by an increase in current density, the morphology of Ni-W tendsto become needle-shaped [100].

4.4.1. Micro/nanocompositeCurrent density plays a key role in the formation of nanocomposite

coatings as well as the amount of incorporated nanoparticles withinthe coating. Regarding incorporation of nanoparticles within Ni-Walloymatrix, it was reported that increasing current density up to an op-timum level consistently leads to increase in nanoparticles contentwithin the coating. However, a further increase in current densityshows an inverse trend. Wang et al. [101] reported that by an increasein cathodic current density up to 10 A dm−2, the embedded diamondnanoparticles contentwere increasedwithin Ni-W coating and droppedthereafter. They also stated that by increasing the cathodic current den-sity (in W-rich electrolyte), tungsten content was enhanced. Aal et al.[102] showed that increasing the current density from 1 to 6 A dm−2

leads to the higher content of SiC nanoparticles and partial increase intungsten and phosphorous contents in Ni-W-P-SiC coatings. It wasdemonstrated that by increasing the current density from 6 to 12 Adm−2, the contents of zirconia and silicon carbide nanoparticles innickel-tungsten coatings are also increased. However, these contentswere diminished when current density exceeded 16 A dm−2 [103,104].

4.5. Rotation rate

Agitation and rotation rates in the bathmay influence both chemicalcomposition and coatings structure [93]. Rotation may change coatingchemical composition by reducing the thickness of diffusion layer overthe cathode surface and increasing the concentration of metallic ions

near the cathode. Diffusion layer over the cathode becomes thinnerwith increasing the rotation rate according to the following equation[105]:

δ ¼ 1:61D1=3ν1=6ω−1=2 ð9Þ

where, δ is diffusion layer thickness, D is diffusion coefficient, ν is kine-matic viscosity, andω is angular rotation rate (s−1). Considering the factthat by increasing the rotation rate in alloy plating processes, boundarylayer thickness is decreased, components with less concentration in theelectrolytewould bemore deposited by themass transport kinetics [49,106]. Younes and Gileadi showed that, independent of bath chemicalcomposition, the mass transfer rate was increased at higher rotationrates; therefore, tungsten contentwas increased in the alloy. Dependingon bath chemical composition, higher rotation rates would lead to in-crease in cathodic current density [49].

Accordingly, when nanoparticles exist in the plating electrolyte in asuspension form, boundary layer thickness is decreased with an increasein rotation rate. Hence, nanoparticles may pass through the boundarylayer more readily and incorporate within the metallic matrix. It was re-ported about diamond nanoparticles that, apart from nanoparticles con-centration in the solution, by an increase in rotation rate by 180 rpm,nanoparticles incorporation is increased and a further increase in rotationrate leads to lowering the amount of incorporateddiamondnanoparticles.Some researchers attributed this fact to collision factor at higher rotationrates [107]. Hashemi et al. [108] studied on SiC nanoparticles in copper-nickel-tungsten coating baths. They reported that rotation rate of400 rpm gives the highest content of SiC nanoparticles within the coat-ings. However, when rotation rate exceeds 400 rpm, a reduction in SiC in-corporation is observed due to excessive flow of the fluid and collision ofun-adsorbed particles. In other works, the optimum rotation rate toachieve the maximum amount of incorporated zirconia and silicon car-bide nanoparticles in nickel-tungsten coatings is stated to be 180 rpm. Itwas explained that fluid flow was not able to bring nanoparticles to thecathode surface at lower rotation rates. Furthermore, less incorporationof nanoparticles at higher rotation rates might cause separation ofadsorbed nanoparticles from the cathode surface [103,104].

On the other hand, grain growth rate is limitedwith diminishing theboundary layer. In comparison, an increase in nucleation rate leads tothe formation of a fine-grain coating. In addition, thinner boundarylayers would cause the deposition of coatings with less surface rough-ness [109,110].

4.6. Different additives

There are many additives applied in Ni-W electroplating baths fordifferent purposes. The most commonly applied additives are sodium

Fig. 5. Schematic representation of pulse current.

985M.H. Allahyarzadeh et al. / Surface & Coatings Technology 307 (2016) 978–1010

saccharin as a stress reliever, sodium bromide for an increase in electri-cal conductivity of the electrolyte, boric acid as a buffer, ammoniumchloride for an increase in current density efficiency, and sodium dode-cyl sulfate and sodium laurel sulfate as surfactants [13]. It was statedthat presence of higher ammonium chloride concentrations led tomore electrochemical deposition of tungsten within the structure[111].Wu et al. [17,112] investigated the effect of 2-butyne-1, 4-diol ad-dition in Ni-W electroplating. On the one hand, this additive causes de-position of very bright and smooth Ni-W films and, on the other, it leadsto a gradual decrease in tungsten content and cathodic current efficien-cy by facilitating hydrogen evolution reaction over the cathode. In thisconnection, 2-butyne-1, 4-diol has been also used as a brightener inNi-W-B electroplating forming coatings with less surface roughness. Inaddition, the optimized usage of this additive may provide fine-grainstructures with high microhardness. Matsui et al. [54] employed 2-butyne-1, 4-diol, sodium saccharin and sodium allysulfonate additivesto improve the tensile ductility of nanocrystalline nickel-tungstenalloy samples. The obtained results showed a 5% improvement in sam-ples deposited in the saccharin-bearing sulfamate bath. In order toachievemore long-term stability in nickel-tungsten electrolytes, an am-monium salt of an organic acid can be used as stabilizer [113]. Sassi et al.[18] used pyridine as an additive in nickel-tungsten electrodepositionfor its leveling effect and corrosion inhibition properties. In the presenceof pyridine, the obtained coatings were a fine-grain, compact, and high-ly corrosion resistant and have a crystalline structure with (311), (111),and (200) preferential planes. The addition of 100 ppm salicylaldehydetoNi-Welectrolytes is reported to cause deposition offine grains aswellas reduction of surface roughness. Kumar et al. [114] showed throughEIS and FT-IR tests that adsorbed salicylaldehyde on nickel-tungstencoatings improves corrosion resistance by diminishing corrosion cur-rent and shifting corrosion potential toward nobler values. Most well-known additives alongwith their roles in nickel-tungsten electroplatingelectrolytes are presented in Table 2.

5. Techniques for Ni-W electrodeposition

Among different deposition techniques available for nickel-tungstenelectroplating, direct current, pulse current, and pulse reverse currenthave been mostly employed. Application of different techniques notonly affects the chemical composition of coatings but also is influentialon their structural properties and roughness features. For instance, puls-ing current (Fig. 5) affects mass and charge transfer during the platingprocess. Anodic current (reverse) may significantly alter structuralproperties and chemical composition of coatings [118,119]. Except theaforementioned techniques, application of ultrasonicwaveswas report-ed to improve Ni-W electrodeposition, and particularly, electroplatingof composite coatings. Indeed, ultrasonic waves are means of hinderingagglomeration of nanoparticles through deposition electrolytes and en-hancing better distribution of nanoparticles within metal matrix [120].For example, it was reported that ultrasonic waves led to the improved

Table 2A brief summary of different additives which applied for distinct proposes in Ni-W electrodepo

Additives Chemical formula Functions

SDS NaC12H25SO4 SurfactantSodium lauryl sulfate C12H25NaSO4 SurfactantSodium saccharin C7H4NNaO3S Stress reliSodium bromide NaBr ConductivBoric acid HBO3 BufferAmmonium chloride NH4Cl To increas2-Butyne-1, 4-diol C4H6O2 BrighteneSodium allyl sulfonate C3H5NaO3S Tensile duDimethyl sulfoxide C2H6OSPyridine C5H5N Leveling aCetyltrimethyl ammonium bromide C19H42BrN SurfactantSodium dodecyl benzene sulphonate, C18H29NaO3S SurfactantSalicylaldehyde C7H6O2 Grain refin

distribution of ZrO2 nanoparticles over the nickel-tungsten matrix andat the frequency of 35 kHz the highest amount of embedded nanoparti-cles were yielded [70]. Ultrasonic waves might be helpful in quality im-provement of Ni-W coatings by making them more compact with lesssurface roughness [16]. Finer grains, less porosity, and reduced internalstresses are mentioned as other advantages of ultrasonic waves [121].

5.1. Direct current technique

In this method, a specific and constant current is applied during thewhole period of the coating process. One of the disadvantages of thistechnique is the creation of residual tensile stresseswhich consequentlyleads to the initiation of cracks within the plated coating. There aremany published data around comparisonsmade between cathodic cur-rent efficiency in direct current and pulse current techniques. Directcurrent efficiency was 61.26% in the bath with gluconate as acomplexing agent while pulse current efficiency in the same bath wasmerely 50% [55].

5.2. Pulse current technique

Pulse technique is an effectivemethod to controlmicrostructure andchemical composition of plated coatings because of its capability in im-proving current distribution and mass transfer processes. Moreover,this technique obviates several problems such as hydrogen evolution,the formation of metallic hydrides, uneven deposits and local pH varia-tions [122–124].

Application of pulse current technique leads to fine-grain nickel-tungsten coatings. It increases tungsten content within the coating [88];although, this matter depends on tungstate or nickel ions concentrationsin the solution [77,125]. It was also reported that coatings became uni-form and hard with improved corrosion properties by employing pulsecurrent technique [126]. In nickel-tungsten electrodeposition, particular-ly when rotating disk electrode is used, electrochemical deposition ofNi2+ ions is kinetically controlled by activation. In comparison, diffusionis the controlling step in electrochemical codeposition of tungstate ions

sition.

References

[13][13]

ever, tensile ductility improvement [13,54,115]ity [13]

[13]e current efficiency [13]r, grain refinement and tensile ductility improvement [17,54,112]ctility improvement [54]

[116]gent, corrosion inhibitor [18]

[102][117]

ement, smoother, adsorbed additive for corrosion resistance [114]

Fig. 6. Cross-sectional SEMmicrographs of Ni-W samples with the same tungsten contentand grain size: (a) conventional direct current at 45 °C, (b) pulse reverse current at 75 °Cand reverse current density of 15 A dm−2. Considerable defects are observed in image(a) while pulse reverse techniques provide samples with premium quality [144].

986 M.H. Allahyarzadeh et al. / Surface & Coatings Technology 307 (2016) 978–1010

[47,127]. Hence, it may be noted that a change in pulse parameters mayaffect kinetic parameters and consequently, leading to different proper-ties of nickel-tungsten alloy and micro/nanocomposite coatings:

5.2.1. Pulse duty cycleDuty cycle is defined according to the following equation [128]:

D:C: ¼ TON

TON þ TOFFð10Þ

It was reported that by a decrease in duty cycle, partial pulse currentdensity for nickel and tungsten was raised which may cause improve-ment in cathodic current efficiency [47,99]. Duty cycle is sometimes in-troduced as the dominant parameter affecting tungsten content ofcoatings. Hou et al. [99] showed that tungsten content and correspond-ingly microhardness were improved by an increase in duty cycle in W-rich bath. They also reported that the major content of incorporatednano-alumina particles was achieved at low duty cycles.

5.2.1.1. Micro/nanocomposite.Duty cycle is considered as an efficient pa-rameter which affects chemical composition, structural features, andsurface properties of Ni-W alloy and micro/nanocomposite coatings. Itis worth mentioning that coating at higher duty cycles generally leadsto the production of Ni-W coatings with more compact morphologies;however, in elsewhere, greater microhardness was reported at lowerduty cycles [11]. Generally, there have been variousfindings of the effectof duty cycle on the incorporation of nano/microparticles on nickel-tungsten alloy matrices. Although many works report that more incor-poration is achieved at lower duty cycles, the opposite is also stated insome publications. For instance, as studied conducted on grapheneoxide shows that by an increase in duty cycle, a higher amount of nano-particles incorporation was achieved [129].

5.2.2. Pulse current densityPulse current density (IP), duty cycle (D.C.), and average current den-

sity (Iave.) are related based on the following equation [128]:

Iave: ¼ Ip � D:C: ð11Þ

One of the advantages of pulse current technique in comparisonwith direct current technique is the possibility to apply higher cathodiccurrent densities. Variations in pulse current density are relatively sim-ilar to those of current density in direct current technique.

5.2.3. Pulse frequencyPulse frequency is defined based on the following equation [128]:

frequency ¼ 1TON þ TOFF

ð12Þ

Franz et al. [47,130] separately investigated co-deposition of tung-sten with nickel in Ni-rich and W-rich electrolytes. Their experimentalresults revealed that by an increase in frequency, the amount of tung-sten co-deposited with nickel was increased in both baths. This fact isattributed to increasing in unsteady statemass transfer rate of tungstatecaused by an increase in frequency. An increase in frequency leads to anarrower diffusion layer which accelerates nucleation and retards graingrowth rate and, consequently, deposition of a fine-grain coatingwith adense structure and less surface roughness [131,132]. It was shown thatnickel-tungsten coatings with fine crystallites were obtained at higherfrequencies [133,134].

5.2.3.1. Micro/nanocomposite. Frequency variations were repeatedlystudied in fabrication micro/nanocomposite coating. Mostly, a frequen-cy increase was reported to raise incorporation of micro/nanoparticles(suspension form in deposition electrolyte) over the alloy matrix. For

example, Fan et al. [135] stated that by increasing frequency from 10to 1000 Hz, more MWCNTs is incorporated in the nickel-tungsten coat-ing. It was also noted that enhancing frequency from 10 to 1000 Hzleads to the increased incorporated titania [134], alumina [99], andgraphene oxide [129] nanoparticulates in nickel-tungsten coatings.

5.3. Pulse reverse technique

Application of an anodic current in form of a reverse pulse duringnickel-tungsten electrodeposition may distinctively affect quality, prop-erties, and chemical composition of obtained coatings. The main reasonof cracking in nickel-tungsten [136] and nickel-molybdenum [137–139]alloys is residual tensile stresses caused by hydrogen reduction on thecathode surface. Hence, the use of an anodic current during electroplatingmay remove hydrogen from the structure and proceed its oxidation overthe cathode surface [140]. Thus, smooth coatingswith the lowest amountof internal stresses are obtained by reverse current technique due to theremoval of hydrogen during application of reverse current [141,142].During electroplating, due to the establishment of proper reduction con-ditions, hydrogen diffuses through the nickel-tungsten lattice and causeslattice expansion. As hydrogen would not remain within the coating, itleaves the structure based on diffusivity and solubility factors. Therefore,hydrogen departure from nickel-tungsten lattice creates residual tensilestresses leading to contraction and finally crack initiation [121,136,143].The differences in coatings' structure obtained by pulse reverse and directcurrent techniques are well illustrated in Fig. 6. Fig. 6a shows cracks initi-ated by internal stresses of a coatingwith inferior quality andFig. 6b dem-onstrates a coating with superior quality made by pulse reversetechnique and elimination of internal stresses. Chemical compositionand grain size may also be controlled by pulse reverse technique andamount of reverse current (Fig. 7). A further increase in pulse reverse cur-rent density leads to the stripping of cathode surface from tungstenatoms; thus, a coating with low tungsten content would be achieved.Fig. 7 shows that by increasing the reverse current density from 0 to 30A dm−2, tungsten dissolution is increased during the reverse currentand consequently tungsten content plummeted from 22 at% to 5 at% in

Fig. 7. Reverse current effects on chemical composition and grain size of nanocrystallineNi-W alloys. The waveform contains cathodic pulse of 20 mS duration and 20 A dm−2

intensity, followed by a 3 mS reverse pulse [144].

Fig. 9. Electron diffraction recorded for Ni–12 at%W alloy film [147].

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nickel-tungsten films. Based on this illustration, the increase in reversecurrent densitymay lead to an increment in crystallite size or grain coars-ening. This fact is well consistent with decreasing inW content in themi-crostructure [144].

6. Microstructural properties

Microstructural properties of Ni-W alloys mainly depend on thechemical composition and plating conditions [145]. Transmission elec-tron microscopy (TEM) studies demonstrate that only solid solution oftungsten in nickel existed in Ni–7–12 at%Walloywhich is in compliancewith nickel-tungsten phase diagram [146] (Fig. 8).

Lattice parameter of FCC structure in N–12 at%W is 0.360 nmobtain-ed from electron diffraction pattern. Regarding the structure of these al-loys, TEM micrographs (Fig. 9) show a bimodal grain size distribution,consisting of 20 to 200 nm sized grains; with twinning and stackingfaults observed in coarse grains.Moreover, nanometer-sized voids spec-ified in TEM micrographs indicate a considerable stress level duringgrain growth [147]. Tunneling scanning microscopy studies show that

Fig. 8. Nickel-tungsten p

the microstructure of nickel-tungsten coatings consisted of nanofibersnormal to the substrate [148].

Generally, face-centered cubic (FCC) structure is indexed for the as-deposited Ni-W alloys. XRD patterns obtained for Ni-W alloys show anintense peak at 2θ = 43.5° corresponding to (111) crystallographicplane. There are two relatively weak peaks at 2θ=50.9° and 74.8° cor-responding to (200) and (220) planes, respectively, related to the purenickel. The dominant texture of Ni-W alloys in (111) plane, commonly,is controlled by the preferential growth in the direction of (111) plane.The Ni-W diffraction peaks, compared with pure nickel diffraction pat-tern, shifted toward the lower 2θ values. According to JCPDS No. 65–2865, pure nickel peaks occur at 2θ = 44.4°, 51.8°, and 76.4°. The shiftof Ni-W peaks toward lower angles, in comparison with pure nickel,may be attributed to the formation of tungsten solid solution in FCCnickel (α-Ni(W)). Tungsten dissolution in nickel matrix leads to expan-sion of nickel lattice. According to the Ni-W phase diagram (Fig. 8),

hase diagram [146].

988 M.H. Allahyarzadeh et al. / Surface & Coatings Technology 307 (2016) 978–1010

tungstenmay dissolve in nickel by 12.5 at% [55,88,149]. Performedmea-surements based on Scherrer's equation for nickel-tungsten alloysyielded nanocrystalline structure and nanoscaled grains [55]. In somecases, a peak was observed at 2θ = 41.4° which may be attributed tonickel solid solution in tungsten (tungsten with body-centered tetrago-nal) [150]. XRD and XPS results, regarding the abnormal compound ob-served at 2θ = 41.8° which faded by annealing at 550 °C, suggest thatthis compoundmay be associated with non-stoichiometric nickel tung-sten oxide (NiWO4) [57,68,151]. Sridhar et al. [151] reported the forma-tion of Ni4W phase (with an observed peak at 2θ = 50.4°) with body-centered tetragonal structure in alloys deposited at extremely low cur-rent densities and high temperatures. When the tungsten contentreaches the critical value of 20 at%, the crystalline structure is changedto amorphous state [150,152,153]. Fig. 10 illustrates the XRD spectrumof a broad range of tungsten contents (from 1.65 at% to 24.2 at%).

Itmay be seen that by an increase in tungsten content grains becomefiner and nanocrystalline structure is converted to an amorphous state.It was also reported that atomic distance, lattice parameter, and fractionof grain boundaries in nanocrystalline Ni-W raise by enhancing in tung-sten content [14,149,154–157]. Herein, Juskenas et al. [158] developed atheoretical linear equation between lattice parameter and tungstencontent:

XW ¼ −7:5208þ 2:13429� αNi−W ð13Þ

where, XW is tungsten atomic percentage and αNi-W is FCC lattice pa-rameter in γ phase. Yoshihisa Kimoto et al. [159] reported a structurecalled amorphous/nanocrystalline duplex composite (Fig. 11) bymeans of nanobeamdiffraction patterns and EDS analysis; this structureconsists of W-depleted and W-enriched phases. Similar results regard-ing W-rich and Ni–rich zone were later reported by Arganaraz et al.[160]. In another research, applying on a wider range of nickel-tungsten alloys, structure, and morphology of coatings with 7 to67 at% tungsten were analyzed by XRD and STM/AFM, respectively.The results showed that by an increase in tungsten content, the nano-crystalline structure tends to become amorphous. Zhu et al. [153] re-ported that by a further rise in tungsten content by 40 at%, thestructurewould undergo another change and is changed to orthorhom-bic crystalline structure (Fig. 12). STMandAFM results showed themor-phology of coatings with a crystalline orthorhombic structure with highdensity regular fine grains, while grains density was also increased athigh tungsten contents. Considering the effect of different fabricationmethods on nickel-tungsten structure, Nasu et al. [161] compared me-chanical alloying with electrochemical deposition techniques. Thestructure of mechanically alloyed nickel-tungsten was reported similarto that of as-deposited; the only difference was the existence of slightcontent of non-dissolved pure nickel and tungsten in the structure ofmechanically alloyed nickel-tungsten. It was reported about Ni-W-Feternary system that by an increase in tungsten content, the crystallinestructure is converted to amorphous and lattice parameter is reduced[162,163].

Fig. 10. XRD spectrum related to the Ni-W electrodeposits with differ

7. Mechanical properties

7.1. Microhardness and Hall-Petch equation

It was shown that amorphous structure patterns appeared by in-creasing the tungsten content by around 20 at%. The further rise in Wcontent up toNi–25 at%Walloy, the structure became completely amor-phous. By the increase in tungsten content and reduction of grain sizedown to 10 nm, the Hall-Petch mechanism is followed and the maxi-mum hardness is obtained at 10 nm grain size. However, hardnessdoes not follow Hall-Petch equation and decreases with finer grains.Yamasaki et al. [61] stated that the reduction of hardness was due to asignificant increase in inter-crystalline volume fraction especially inareas related to triple junction zone. On the other hand, Hall-Petchbreakdown regime was reported to occur at a grain size of 14 nm and8 nm for nominally pure nickel and plated nanocrystalline nickel, re-spectively. When alloying element (tungsten solid solution) slowly dif-fuses in nickel, the inverse of Hall-Petch equation might be satisfied inthe finest grains according to diffusive creep mechanism and grainboundary sliding theory [66,164]. In this regard, Schuh et al. [144,165]expressed that production of pure nickel with a grain size lower than10 nm is difficult; however, this is feasible through alloying as for Ni-13 at%W with a grain size lower than 10 nm, which demonstrates theimproved hardness and wear resistance. It is seen in Fig. 13 that by in-creasing tungsten content from 16.9 at% to 18.8 at%, hardness is raisedfrom 650 to 780 Hv. The further increase in tungsten content by24.2 at% reduced the hardness to 550 Hv. With nanoscaled grain size,these changeswere not due to the strengthening effect of the solid solu-tion; yet, they were caused by the Hall–Petch and the reverse Hall–Petch effects [90].

Considering the fact that grain size is reduced by an increase in tung-sten content in nickel-tungsten coatings [13,14,166], the transition be-tween the classical and Hall-Petch behavior which can be quantitativelyintroduced by developing the critical rain size. In this regard, Schuhet al. [66] quantitatively defined the critical grain size (diNi-W) in transi-tion region as the following formula:

diNi−W ≈DW

DNi1−cð Þ þ c

� �2=7diNi ð14Þ

Assuming DW/DNi ≈ 5×10−5 at room temperature, for alloy Ni–13 at%W, this calculation roughly gives dNi−W

i /dNii ≈ 0.56. This valuesuggests that the inflection grain size would be reduced by about a fac-tor of two by alloying Ni with 13%W. This equation calculates approxi-mately the critical grain size at a specific chemical composition inwhichthe Hall-Petch breakdown occurs. It is assumed that emergence ofboundary diffusive processes leads to Hall-Petch breakdown. InEq. (14), diNi is the critical grain size of pure nickel, DW andDNi are nickeldiffusion coefficients in tungsten and nickel, respectively, and c is tung-sten to nickel atomic ratio. Giga et al. [94] investigated the Hall-Petchbreakdown in nanocrystalline nickel-tungsten alloys electroplatedwith grain sizes of 5, 8, 12, and 20 nm by means of tensile testing. It

ent amount of tungsten from 5 wt% to 50 wt% a [14] and b [154].

Fig. 11. Tungsten concentration visualizations on cross section of the Ni-W alloys with grain size of (a) 5 nm and (b) 8 nm. The regions with low concentration of W are painted in darkcolor (blue), while high tungsten concentration are shown in brighter color (light green) [159].

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may be observed in Fig. 14 that at the critical grain size of 8 nm, the clas-sical Hall-Petch is converted to its inverse form. It was also reported thatby reducing grain size from 20 to 8 nm, both hardness and tensilestrength are enhanced. Contrarily, hardness and tensile strength werediminished with further decrease in grain size below 8 nm. Sriramanet al. [62,67] studied the dependence of hardness and slidingwear resis-tance on Hall-Petch for coatings electrodeposited at 75 °C and 85 °C.Their results revealed that hardness and wear resistance are improvedby reducing grain size to 20 nm at 75 °C. However, wear resistancewas degraded at lower grain sizes (below 20 nm) due to the brittle frac-ture. Hall-Petch breakdown was reported for alloys electrodeposited at85 °C,where the bestwear resistance corresponded to the coatingswith6 to 8 at%W (with a crystallite size of about 15–20 nm).

Rupert et al. [167] studied abrasive wear response of nanocrystallineNi-W alloys (5–105 nm) and its dependence on Hall-Petch breakdown.They stated that with coarse grains (8–105 nm), wear resistance signif-icantly was improved by reduction of grain size. The wear resistance ofthe finest grain size (5 nm) specimen is found to be higher than thatpredicted based on hardness alone. This deviation from Archard scalingor Hall-Petch breakdown is due to mechanically-driven structural evo-lution which occurs during wear, including grain growth and grain

Fig. 12. X-ray diffraction patterns related to the electrodeposited Ni-W alloy with 7–67 at%W [153].

boundary relaxation. Rupert and Shun works proved that microhard-ness of nanocrystalline Ni-W films did not follow Hall-Petch Eq. (15)for all grain sizes:

H ¼ H0 þ kd−n ð15Þ

where, H is themeasured hardness, H0 is the hardness of a single crystal,k is material's constant, and d is the average grain size. In Hall-Petchequation, n is 0.5. Regarding themicrohardness of nanocrystalline coat-ingswith coarser grains, Hall-Petch is well satisfied. For instance,micro-hardness of coatings with grain sizes between 15 and 47 nm properlyobeys the Hall-Petch. However, microhardness values of samples withgrain sizes varying from 3 to 10 nm did not follow Hall-Petch (Fig. 15)[166]. Similar results were obtained in other works so that the highesthardness of electrodeposited nickel-tungsten alloys is obtained in coat-ings with grain sizes within the 10–15 nm range [90,168]. Althoughhardness, grain size, and structure of nickel-tungsten coatings stronglydepend on tungsten content, they might be also affected by coating/substrate interface in extremely thin Ni-W coatings. This behavior isdue to the preferential formation of a Ni-rich film near the substrate be-yond which Ni-W alloy is deposited. The Ni-rich film is designated as‘intelligent interphase region’ by Pisarek et al. [169].

Considering into account that grain size straightly affects nickel-tungsten alloys, a great attention is paid to parameters that make alloygrains finer and consequently improve hardness. It was reported thatpresence of nanoparticles such as alumina [111] leads to the creationoffine grain structures due to increase in nucleation sites on the cathodesurface [170]. Therefore, coatings hardness may be enhanced by forma-tion of nano/microcomposites incorporating various kinds of nano/mi-croparticles like TiO2 [126,134], TiN [116,171], ZrO2 [103], Ce3F [172],W [173], diamond [101,123], Si3N4 [174], BN [175], and SiC [176] inNi-W alloymatrix bymeans of different mechanisms such as dispersionstrengthening and particle strengthening. It was reported that coatingscontaining 36.2 at% diamond nanoparticles with a hardness of 2249 ±23 Hv were electrodeposited by vertically positioning of the cathodein the plating electrolyte. The hardness of these coatings can be en-hanced up to 2647 ± 25 Hv by heat treatment of these coatings (at600 °C for 1 h at Ar atmosphere) [177].

7.2. Tensile strength

Electrodeposited nickel-tungsten alloys with nanocrystalline struc-ture generally possess proper ductility and tensile strength. However,the presence of hydrogen on cathode during electrodeposition severelydiminishes their strength. It is possible to achieve a tensile strength of2333MPa as well as high ductility and complete bending to 180° (with-out fracture) in Ni–20.7 at%W alloys with average grain size of 3 nm

Fig. 13.Microhardness as functions of tungsten content (a) and D−0.5 (b) of the Ni-W electrodeposits [90].

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through degasification [152]. Yet, tensile strength for Ni-W alloyswith agrain size of 8 nm was reported in the range of 800–1050 MPa. The in-accuracy in tensile strength values was attributed to the presence ofnanopores through the microstructure [178]. Elsewhere, tensile testswere carried out for nickel-tungsten alloys with grain sizes of 2, 5, 8,12, and 20 nmby Giga et al. [94]. They stated that the increase in tensilestrength continued with decreasing in grain size to 8 nm and the max-imum tensile strengthwas achievedby1238MPa and itwas diminishedatfiner grains (2 and 5 nm) because of theHall-Petch breakdown. Nom-inal tensile strength and total elongation at fracture for monolithicnanocrystals Ni–16 at%W with 10–20 nm grain size were 1700 MPaand 1%, respectively. Ymasaki et al. [179] produced some micrometer-sized holes in nickel-tungsten alloys by means of UV-lithography tech-nique. They showed that with finely dispersed array through-holeswhichwere 30 μm in diameter, nominal tensile strength and elongationreached 2200 MPa and 1.2%, respectively. Several parameters of tensiletest for array through-holes and monolithic samples are presented inFig. 16. It may be perceived from that stress concentration factor (α)which is induced by array through-holes increased by a maximum2.14 for 20 μm array through-holes. Thus, it reduced to 1.65 by a furtherincrease in diameter by 50 μm after formation of array through-holes;apparent yield stresswas decreased by 46% and 60% to that ofmonolith-ic samples with diameters of 20 μm and 50 μm, respectively. Arraythrough-holesmay lead to accumulation of shear bands and cracks (ini-tiated from array through-holes) in Ni-W alloys which, in turn, en-hances mechanical strength and ductility under tensile test conditions.

Fig. 14. Stress–strain curves of the electrodeposited Ni–Walloys. The corresponding grainsizes are denoted in each graph [94].

Fig. 17 illustrates a stress-strain curve of the nickel-tungsten samplewith an amorphous structure (with the minimum fluctuation in chem-ical composition) produced by electroplating on a copper substrate. Thetensile strength of amorphous nickel-tungsten samples was reported tobe 2681 MPa with a standard deviation of 6.1%. Amorphous nickel-tungsten alloys demonstrate a wide 1–2% elastic area and 1–3% plasticstrain. It was reported that elastic limit of 2% is in good agreementwith cast amorphous alloys. However, the compressive strength of2.7 GPa is relatively enormous as compared with that of other amor-phous materials [180].

7.2.1. Residual stresses in Ni-W electrodepositsFormation and presence of residual stress severely challenge the

electroplated nickel-tungsten alloys. Matsui et al. [53] made an attemptto improve the tensile ductility of deposited nanocrystalline Ni-W fromsulfamate bath using propionic acid. They managed to produce nano-crystalline Ni–0.5 at%W with a tensile strength of 1.3 GPa and tensileductility of 2.7%. In addition, Ziebell and Schuh [63] investigated nano-crystalline Ni-W in grain size range of 4–63 nm and thickness of 10–100 μm. Based on their findings, apparently, twomajor factors associat-ed residual stresses are: 1) the excess free volume in grain boundariesthat originates fromprocess conditions, and 2) the total volume fractionof grain boundaries which is controlled by grain size. Tensile strength is300–2300MPa in these alloyswith the absolutemaximumbelonging tothe sample with grain size of 15 nm. Nakayama et al. [181] developed anovel agitation method named “brushing technique” for fabrication of

Fig. 15. Hardness of nanocrystalline Ni-W versus grain size (in logarithmic scale). Thisgraph shows that only coarser grains follow the Hall-Petch (dashed line) [166].

Fig. 16. Studying of different tensile test parameters for nanocrystalline Ni-Wwith finely dispersed array through-holes with micrometer size [179].

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amorphous and nanocrystalline nickel-tungsten alloyswith high tensilestrength (about 2500 MPa) and a large plastic strain of 2.9% with workhardening. In this technique, hydrogen bubbles are directly removedfrom the cathode surface. It was reported that nanocrystalline nickel-tungsten alloys with high tensile ductility (about 5%) could be obtainedby addition of saccharin to sulfamate bath. Besides, texture control playsa significant role in the improvement of tensile ductility of nanocrystal-line Ni-W alloys. Hence, the dominant texture in samples with highplastic deformability is (200) while the dominant texture in specimenswithout plastic deformability is (111) [54]. Considering the lack of nomeaningful relationship between tensile ductility and grain size (ortungsten content), tensile behavior using orientation index for (200)plane may be explained as follows: When orientation index is lessthan unity, no plastic deformationmay be observedwhile tensile ductil-ity increases at higher indices. Thus, the strength of (200) texture has tobe considered as a requirement for tensile ductility of nanocrystallineNi-W alloys [182,183]. Yamasaki et al. [184] reported the tensilestrength of Ni–14–24 at%W alloys to be about 3 GPa. As Fig. 18 shows,nominal tensile strength of nanocrystalline Ni–14.4 at%Wwas obtainedto be 2.9 GPa. A tensile strength exceeding 3 GPa was obtained with anincrease in tungsten atomic content by 18.2% while it was dropped to2.9 GPa again at tungsten content of 23.6 at%.

Fig. 17. Compressive stress-strain curve of bulk amorphous Ni-W alloy [180].

It is recently reported that addition of 5 vol% of SiC in Ni-W electro-deposits enhanced the coating hardness and elastic modulus, respec-tively, from 10.31 ± 0.65 GPa to 14.32 ± 0.63 GPa and from 119.74 ±3.15 GPa to 139.26 ± 2.09 GPa. These increments directly lead tostrengthening in the coating. The boosted internal strain and grain re-finement simultaneously increased the dislocation density, which inturn promoted the formation of grain boundary misorientations andnano twinning [176]. It is reported that incorporation of ZrO2 particlescould influence Youngmodulus. In addition, by the optimum incorpora-tion amount of zirconia (about 1.5 at%) the highest Yang modulus(205 GPa) and hardness (9.1 GPa) are achieved [185].

7.3. Creep and fatigue

The steady-state creepbehavior of Ni–0.6, 1.3, 2 at%Wsolid solutionswas studied using constant stress tensile creep tests in the temperatureand stress range of, respectively, 850 to 1050 °C and in 20.7 MPa to48.3 MPa in a vacuum. After correcting temperature dependence ofthe elastic modulus, activation energies for creep were observed to be71.4±2.0, 74.4±3.0, and 75.8±2.0 kcal permole for alloys containing0.6, 1.3, 2 at%W, respectively. The steady-state creep rates demonstratea power law stress dependence with an exponent, n, equal to 4.8 ± 0.2for all alloys studied [186]. This value is about 3 for Ni-3.5 at%W inOikawa et al. research [187], where the steady-state creep of Ni-3.5 at%W alloys was investigated in the temperature range of 750 to1000 °C and at the stress level from 2.0 to 9.0 kg/mm2. The creepcurve shows often three stages including normal transient, steady-

Fig. 18. Stress-strain curve from tensile test of nanocrystalline and amorphouselectrodeposited Ni-W [184].

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state, and accelerating; through which the value of internal stress de-creases with increasing at creep temperature and increases with in-creasing applied stress. Matlock and Nix investigated the effects ofhelium on the high-temperature creep properties of Ni-2 at%W. Theyreported that the presence of helium bubbles at grain boundaries pro-motes premature fracture early in the primary creep stage by decreas-ing the effective fracture surface energy and therefore inhibits theobservation of steady-state creep [188]. In another research, creep testresults showed that the elongation strain increases with increasing theapplied tensile stress at 300 °C or lower temperatures attaining to amaximum value of 0.68%. The elongation strain tended to saturate ateach testing temperature and the size of nickel-tungsten nanocrystalsreached 10–15 nm. It is probable that applying high tensile stressesmight be effective in the uniform plastic deformation of nanocrystallineNi-W alloys [189]. Regarding the effect of grain size on fatigue life ofelectrodeposited nanocrystalline Ni-W alloys on carbon steel substrate,Sriraman et al. [62] stated that non-coated samples showed a superiorfatigue life comparedwith coated samples. Such a prolonged lifewas at-tributed to inherent microcracks and tensile residual stresses in coatedsample which degrade fatigue life. However, fatigue life was reportedto increase with a decreasing grain size in comparisons made amongcoated samples. In a similar study, Mroz et al. [190] showed that fatiguelife of electrodepositedNi-Walloys (on steel substrates)was dependenton stress amplitude. In stress-controlled fatigue experiments, coatedsamples showed a lower resistance against fatigue than non-coatedsamples at high-stress amplitudes. However, the fatigue strength ofboth samples was of the same value at low-stress amplitudes. Indenta-tion techniques illustrated a cyclic softening as a result of fatigue dam-age. Moreover, damage process of coatings initiated around a surfacedefect on the center of samples loaded with 400 MPa ≤ σmax; for sam-ples loaded with 400 MPa ≥ σmax, fatigue was only observed with dom-inant crack propagation on the edge of samples and no significantchange occurred on the surface until the final fracture. Generally, it isworth mentioning that tungsten content may be of great impact on fa-tigue behavior of nickel-tungsten alloys so that fatigue resistance re-duces with an increasing tungsten content in spite of some increase inhardness. This fact is due to the increase in internal residual stresses.The decrease in fatigue life as a result of tungsten increase may be alsoattributed to a reduction in grain size and theory of grain boundariessliding [66,191].

7.4. Fracture toughness

Fracture toughness of nanocrystalline nickel-tungsten coatings withfunctionally graded structure (with an increase in tungsten content to-ward the surface) was investigated by Wang et al. [192]. Toughness ofNi-Wwas decreasedmoving from the substrate toward the coating sur-face by an increase in tungsten content.

Fig. 19. The fracture cross sections of the funct

As shown in Fig. 19, fracture cross-section of FG nickel-tungstencoating shows that some splitting fracture ribs are formed with goodplastic deformation in direction with the crystalline columnar growthwhich indicates high toughness inside the coating. By an increase incoating thickness, the height of splitting fracture ribs declined and brit-tle cleavage was stored in the middle of fracture. In other words, therewas a mixture consisting of ductile and brittle cleavages in the middleof fracture while river tracks and several cracks are visible adjacent tothe outer fracture zone (near to the coating surface) that implies classicfracture [192]. Many observations have been made on annealing tem-perature effect on fracture toughness behavior without any unique con-clusion. As an explanation, it can be stated that annealing of nickel-tungsten coatings at higher temperatures does not change fracturetoughness in a vivid and simple way and any change in fracture behav-ior may depend on the nature of microstructure and existing inclusions[193].

Considering brittleness and assuming the linear elastic behavior ofNi–17.5 at%W alloy, fracture toughness values were obtained in therange of 1.49–5.14 MPa m0.5. Accordingly, fracture toughness of Ni–12.7 at%W alloy was measured to be in the same range. These valuesare higher than those of silicon and silicon-based alloys (0.83–0.95MPam0.5) but lower than those obtained for nickel and nickel based al-loys. The average fracture toughness for Ni–17.5 at%W was reported tobe 2.97MPam0.5 so that this alloymay be nominated forMEMS becauseof its proper mechanical properties [39,194]. Since local flaking ofnickel-tungsten coatings is mainly attributed to the high brittleness ofthese coatings, this issue may be resolved to some extent by improvingthe fracture toughness of coatings and strengthening its bonding withthe substrate through modification and optimization of electrodeposi-tion conditions [195]. Furthermore, cracks initiation in nickel-tungstencoatingswith an amorphous structure ismostly due to residual stresses;hence, the addition of some alloying element like iron [65]may improvethe coating toughness and adhesion.

8. Magnetic properties

The structural origin of perpendicular magnetic anisotropy (PMA) inelectroplated nickel-tungsten thin films was investigated by NicolaSulitanu [196]. Perpendicular magnetic anisotropy and magnetizationdeviation were found in transition region with tungsten content of 4–5 at%W due to the presence of tungsten and probably compositionalinhomogeneity. Saturation magnetization, magnetic anisotropy field,perpendicular magnetic anisotropy, and perpendicular coercivity of Ni–4.5 at%W were reported as 420 kA m−1, 451 kA m−1, 230 kj m−1 and113 kA m−1, respectively. Considering the tendency of tungsten atomsfor segregating grain boundaries, driving mechanism was announcedto be required for the appearance of PMA in the interface of themagneticcore of nickel crystalline grains and the Ni–W non-magnetic boundary

ionally graded Ni–W electrodeposit [192].

Fig. 21. Magnetization curves (a) for deposited Ni (with 10 mm thickness) and(b) electrodeposited Ni–15 at%W alloy with a 3 mm thickness. Curve (c) relates to Ninanowires with 80 nm diameters and 6 mm lengths and (d) illustrates Ni–15 at%W alloynanowires with 80 nm diameters and 6 mm lengths. Dashed lines: magnetic field wasapplied in-plane. Solid lines: magnetic field was applied perpendicular to the film's plane[198].

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layer. Magnetoelastic anisotropy associated with in-plane internalstresses and positive magnetostriction plays a vital role in the appear-ance of PMA. Also, tungsten contentmaybe influential in control of inter-nal stresses magnetization. Magnetocrystalline anisotropy of b111Ncolumnar grains and their shape magnetic anisotropy can be introducedas another source of PMA. Besides, there is a direct connection betweenin-plane internal stresses and grain columnar growth in nickel-tungsten alloy films. It is well known that grain size decreases with in-creasing tungsten content but lattice constant still increases. As shownin Fig. 20, magnetic coercivity force in nickel-tungsten alloys wasdecreased to 10 Oe at high tungsten contents considering thatmagnetocrystalline anisotropy reduces with a grain size decreasing[197].

As shown in Fig. 21, isotropic magnetization response of electrode-posited nanocrystalline Ni-W alloy nanowire arrays reveals that magne-tization of Ni–15 at%W thin films is saturated at 2.5 kOe in aperpendicular direction to the plane, whereas the thin film of pure nickelis hardly magnetized in the perpendicular direction. Contrarily, Ni–15 at%W alloy nanowire arrays easily saturated at 1.0 kOe even in a per-pendicular direction to the array film plane because of the long-axis di-rection of alloy nanowires. Effective magnetic field (Heff) in the alloyarrays is expressed via the following equation:

Heff ¼ Ha−Hd ¼ Ha− f d �M=μ0 ð16Þ

where, Ha is the applied magnetic field, Hd is the demagnetizing field, fdis a factor of the demagnetizing field, M is a magnetic moment and μ0 ispermeability constant. According to Fig. 21 and Eq. (16), one may con-clude that if the appliedmagnetic field was in a perpendicular directionto the plane, fd and Hd would become maximum; hence, Heff is mini-mized. Also, fd and Hd are minimum in the in-plane direction and Heff

is maximized and reaches Ha [198].Esther et al. [163] tried to achieve an optimumcomposition in Ni-Fe-

W alloys with desired magnetic properties (with high tungsten con-tent) by studyingmagnetic properties of nickel-tungsten thin films con-taining iron. It was revealed that coercivity (Hc) and saturationmagnetization were increased and decreased, respectively, with an in-crease in tungsten content. In this regard, Ni62Fe25W14 was reportedas the optimum composition of alloy films with soft magnetic behavior(low coercivity and high magnetization). The coercivity and saturationmagnetization of these alloys are 8 Oe and 0.99 T, respectively. It wasalso seen that saturation magnetic flux density (Bs) and coercivity de-creased and increased, respectively, with an increase in citrate

Fig. 20. Relationship between W, Cr, and Mo in alloy with magnetic coercivity [197].

concentration in the electrodeposition solution. In fact, citrate additionaffects magnetic saturation of alloy by dilution mechanismwith chang-ing tungsten content as a non-magnetic alloying element.

Fig. 22.Weight gain vs. oxidation time curves at 700 and 800 °C in air for three Ni-W andbulk Ni specimens [200].

Fig. 24. Thickness of inter-diffusion layer after oxidation at 600 and 700 °C for depositedNi-W (on steel substrate) [204].

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9. Oxidation and high-temperature corrosion

Oxidation behavior of Ni–16 at%W coatings (deposited on STD61carbon steel substrates) at 700 and 800 °C was investigated by Leeet al. [199,200]. Fig. 22 illustrates and compares the weight gain curveversus oxidation time at 700 and 800 °C for three Ni-W alloys andbulk Ni. Studies showed that specimens consisting of nanocrystallinegrains supersaturated with tungsten, molybdenum, and silicon with alattice parameter of 7.782 Å prior oxidation. There are four distin-guished regions in Ni-W coatings after oxidation tests: the first regionconsists of NiO outer layer and the second region is made mixed oxidesinner layer (NiO+NiWO4), which the outer layer indicates the upwarddiffusion of nickelwhile the inner layer corresponds to downward diffu-sion of oxygen. The third region is the unoxidized region of coating andexactly placed beneath the oxide layers (the first two regions). Nickelsolid solution supersaturated with O, Fe, andW is the main constituentin the third region. Ni-W-Fe compounds were formed in the fourth re-gion exactly above the substrate/coating interface. These Ni-W-Fe pre-cipitates are surrounded by Fe-enriched, Cr-containing Ni grains. SEMcross-sectional micrographs with EPMA analysis and schematically oxi-dation mechanism of nickel-tungsten coatings at 700 °C and 800 °C areillustrated in Fig. 23a, b, and 23c. However, regarding the formation ofternary NiWO4, it was also reported that below 1000 °C no evidenceof NiWO4was found in Ni-3 at%W alloy due to the lowmobility of tung-sten which prevents chemical reactions [201]. The oxidation kinetic forNi-4.9 at%W in air at 1000 °C exhibits a parabolic oxidation rate, whichindicates a diffusive behavior of species during the oxide formation. Themulti-layered structure of the formed oxide consists a pure nickel oxideouter layer (NiO), a porous NiOmatrix containingNiWO4 second phase,and an internal oxidation zone with sub-micron WO3 oxides which isconverted, progressively, to NiWO4 precipitates in the alloy [202]. Ni-15 at%W cannot be a potential candidate as a sulfidation-resistant coat-ing. In the SO2 atmosphere, this coating is corroded severely due to thepresence of S and formation of less-protective NiO, NiS, WO3, andNiWO4; however, most of the scale consisted mainly of not sulfidesbut oxides [203].

Traces of Fe2O3 andNiFe2O4 onNiOdemonstrate upward diffusion ofiron. Besides, nickel also partially is diffused into the substrate. Weight

Fig. 23. Cross-sectional micrographs (a) and EPMA analysis of A–B (b) of electrodeposited Nmechanism of Ni-W coatings at 700–800 °C [199].

gain evaluations show that Ni-W oxidation kinetics follows the parabol-ic rate law. In addition, Ni-Wcoatings exhibit an inferior corrosion resis-tance compared to pure nickel at 700 °C and 800 °C probably due to theformation of NiWO4 in coatings. Eraslan and Ürgen studied oxidation ofNi-W-B alloys at 600 °C and 700 °C. Experimental results revealed thatNi-W-B alloy ismore resistant to corrosion than bothNi-B andN-P coat-ings (Fig. 24). Although all samples display a rather same oxidation be-havior, it is stated that tungsten had a distinctive role in the reduction ofinter-diffusion layer thickness. The presence of tungsten (about0.65 at%) develops a barrier which delays iron diffusion into the coatingand consequently the formation of iron oxides on the coating surface[204]. It is reported that La2O3 nanoparticles improve oxidation andhigh-temperature corrosion resistance of Ni-W coatings. Oxidation re-sistance of Ni-W coatings at 700 °C was improved with an increase inLa2O3 nanoparticle content from 3% to 13%. In addition to oxide disper-sion strengthening, La2O3 nanoparticles reduce the interface betweencoating and air and provide a strong inhibition effect on the diffusionof oxygen by doping with coatings [205]. It is reported that CeF3 nano-particles improved oxidation resistance of nickel-tungsten coatings at

i–16%W (on steel substrate) after oxidation at 800 °C for 5 h and schematic oxidation

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650 °C. Weight gain experiments show that weight gain of coatings dueto oxidation decreased by 60% with the incorporation of CeF3 nanopar-ticles up to 1.6-2 at%. The improvement in oxidation resistance is attrib-uted to the formation of a compact and fine structure as well asdispersive strengthening caused by the addition of CeF3 nanoparticles.Chemical neutrality and extremely low electrical conductivity of CeF3nanoparticles which are uniformly dispersed in Ni-W matrix isolatecoating surfaces in an oxidative medium and improve oxidationresistance [172].

10. Corrosion properties

Corrosion behavior investigation of nanocrystalline nickel-tungstencoatings investigated in 3.5% NaCl solution at pH = 3 (adjusted withsulfuric acid) and at pH = 10 (adjusted with NaOH) show that inbasic media coatings are more resistant to corrosion than acidicmedia. Corrosion resistance of nickel-tungsten coatings was reportedto be affected by three parameters: 1) tungsten content, 2) volume ofgrain boundaries which may act as potential sites for corrosion reac-tions, and 3) the crystallographic texture of coatings. It was reportedthat, in basic NaCl solutions, the corrosion rate of coatings increasedwith a decreasing grain size. On the other hand, in acidic NaCl solutions,the corrosion resistance improved with lower grain sizes and passiv-ation degree increased at high tungsten contents [14]. However, it wasreported in the literature that pH values (3, 7 or 10) do not significantlyaffect corrosion behavior of Ni-W coatings in saline media and there isno close relationship between tungsten content and corrosion behavior.Since coatingswith finer grain size are associatedwithmore dissolutionand on the other hand, an increase in tungsten contentmay provoke aninhibitory behavior or a diffusion barrier actionwhich, in turn, improvescorrosion resistance. Thus, the most important factor affecting the cor-rosion resistance of nickel-tungsten coatings is considered to be surfaceand morphological imperfections originating from coatings applicationconditions [150,154]. The effect of chloride (Cl¯) concentration (0.03,0.3 and 1 M) on corrosion behavior of Ni-W coatings was also investi-gated and results showed that concentration variations did not have asignificant effect on corrosion potential and current and theonly param-eter bringing dramatic increase in corrosion rate is temperature. It waspointed out, elsewhere, that tungsten (whichmay result in the prepara-tion ofmore stable films) is themain controlling parameter of corrosioncurrent. Some authors reported that corrosion resistance in nickel-tungsten coatings is lower than that pure nickel samples [14]. It wasstated, in this regard, that nickel-tungsten coatings had amore negativecorrosion potential (more active) which is even shifted toward furthernegativity with an increase in tungsten content. Besides, nickel-tungsten coatings showed higher corrosion current values than purenickel. However, opposing observations have also been reported [92].According to the anodic branch of Ni-W coatings polarization curve, an-odic dissolution occurs at much less speed in these coatings in compar-isonwith pure nickel samples. Herein, Ni-Walloys havemore resistanceagainst localized corrosion (i.e. pitting corrosion) inNaCl solutions com-paredwith pure nickel [150,206]. However, there is a report aboutmoreanodic dissolution and instability of the passive film owing to highercontents of W in the coating [157]. Yang et al. [16] compared corrosionresistance of crystalline, nanocrystalline, and amorphous nickel-tungsten coatings in 0.5MNaCl solutions. Their experimental results re-vealed the higher corrosion resistance for amorphous coatings thancrystalline ones, while nanocrystalline coatings displayed the highestcorrosion resistance of all. The main explanation for this observationmight be the preferential dissolution of nickel and formation of W-rich film over the surface inhibiting further corrosion. For this reason,a passivation behavior is commonly observed in these coatings, apartfrom chemical composition, in the range of 500 to 1500 mV above theOCP in a 1 N sulfuric acid solution [13]. In spite of what mentioned ear-lier about shifting the corrosion potential towardmore active regions asa result of an increase in tungsten content [150], it was also reported

that corrosion resistance of Ni-W coatings in saline solutions is im-proved with increasing the tungsten content up to 7.54 at%. Further in-crease in tungsten content, however, diminished corrosion resistance.This fact may be attributed to the formation of fine grains through thecoating microstructure with increasing tungsten content. In otherwords, an increase in tungsten content initially enhanced the corrosionresistancewhereas a further rise in tungsten content led to grain refine-ment which, in turn, is responsible for the increase in grain boundariesvolume and drop in corrosion resistance. However, with an increase innumber and equal distribution of grains, corrosion does not progress lo-cally and occurs uniformly over the surface [13]. Nevertheless, non-uniform corrosion of nickel-tungsten alloys in NaCl media was also re-ported by some researchers [127]. It was mentioned elsewhere thatthe crack density initiated over the surface is to be considered as themain parameter in corrosion resistance of Ni-W films [15]. ComparingRaman spectra illustrated in Fig. 25a and b revealed that corrosion prod-uctswere strongly affected by acidity and concentration in the corrosivemedium. According to Raman spectra (Fig. 25a), corrosion products ofnickel-tungsten coatings in 0.3 M NaCl solution were mostly tungstenoxide (WO3) and nickel oxide (NiO); while Na(OH)2, NiO, and WO3

were the most prevailing corrosion products in 0.1 M NaCl basic solu-tion. In Raman spectrum, the observed peaks at 367 cm−1, 535 cm−1,770 cm−1 and 960 cm−1 were indexed to Ni-OH, Ni-O, O-W+6-O andW+6 = O bonds, respectively [127,154].

It is reported that annealing treatment may improve the corro-sion resistance of nickel-tungsten coatings; and regarding the heattreated specimens the best corrosion resistance corresponds to thetemperature of 500 °C [207,208]. It was recently shown that corro-sion resistance of nickel-tungsten coatings is enhanced by fabricat-ing multi-layered structures and composition modulated multi-layered alloys [209]. The addition of molybdenum to the binarynickel-tungsten alloys also may improve the corrosion resistance,particularly, at anodic potentials. Electrochemical impedance spec-tra and local pH measurements exhibit that presence of molybde-num leads to the formation of secondary oxide and hydroxide filmswhich restrain further dissolution of coatings [210].

10.1. Micro/nanocomposite

The presence of various nano/microparticles may reduce the corro-sion rate in nickel-tungsten coatings. For instance, incorporation ofMWCNT and the increase in its content shifts the corrosion potential(Ecorr) toward nobler values [211]. Also, corrosion current (icorr) fallsat first reaching a minimum point and then again tends to increase. In-crease in corrosion current due to high contents ofMWCNT is attributedto agglomeration of those particles which leads to degradation of de-sired electrochemical properties of coatings [135,212]. Experimentsshow that the size of nano/microparticles utilized into nickel-tungstenbased nano/microcomposites is of considerable importance in corrosionproperties. Although pitting corrosionwith nanosized poreswas report-ed in several cases, considering the effect of average size of alumina andCNT particles (from micro to nanometer scales), corrosion current wasshown to decrease by 37% after 192 h of immersion with the reducingaverage size of alumina and CNT particles from 724 nm to 9 nm [213].The addition of TiO2 nanoparticles in nickel-tungsten alloy matrixshifted corrosion potential toward nobler values and improved the cor-rosion resistance of Ni-W coatings by enhancing charge transfer resis-tance and reducing corrosion current density [126]. It was stated thatSiO2 nanoparticles improve the corrosion resistance of Ni-W coatingsby acting as physical barriers and filling pores, grooves, and voids overthe surface of nickel-tungsten coatings. SiO2 nanoparticles shift the cor-rosion polarization curve to the left direction and transfer the corrosionpotential toward nobler values. However, it is reported that excess SiO2

nanoparticles leads to degradation of corrosion resistance and shiftingthe curve towardmore active values [51,214]. The presence of SiC nano-particles in nickel-tungsten alloy matrix trebled charge transfer

Fig. 25. Raman spectra of corroded surface of depositedNi-Walloy in 0.3MNaCl [154] (a) and 0.1MNaCl basic solution [127] (b). Thisfigure compares obviously the corrosion products ofNi-W alloys in different media.

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resistance within the coating and improve corrosion resistance byshifting corrosion potential toward nobler amounts [117,215]. TiN[171], BN [175], and ZrO2 [100] nanoparticles are also reported to be ef-fective in achieving nobler corrosion potentials. Titanium nitride andboron nitride particles double the corrosion resistance of nickel-tungsten coatings by increasing charge transfer resistance as well as re-ducing corrosion current density [87]. Silicate addition improves thecorrosion resistance of nickel-tungsten coatings significantly. Incorpo-ration of PTFEor PCTFE particles in nickel-tungsten coatings successfullyshifts the corrosion potential toward nobler values. Increase in PTFE orPCTFE content, also, reduced corrosion current density and consequent-ly enhanced corrosion resistance [72,216]. Uniform and compact coat-ings with proper corrosion resistance can be obtained by addition ofgraphene oxide to nickel-tungsten alloys. Electrochemical measure-ments reveal that graphene oxide improves the corrosion resistance ofnickel-tungsten coatings by shifting corrosion potential toward noblervalues and reducing corrosion current density [129]. Fig. 26 briefly sum-marizes and demonstrates the effect of different particles on a typical E-log(i) curve in chloride solution.

Fig. 26. Brief effects of W, Mo, and different micro/nanoparticle effect on corrosionproperties of Ni-W alloy in chloride solution in a typical polarization curve.

11. Tribological and wear properties

Since introducing nickel-tungsten coatings as proper replacementsfor hard chromium coatings in applications demanding a high hardnessand wear resistance, their wear behavior and properties have been fre-quently subject to different experiments. In this regard, Sriraman et al.[67] studiedwear resistance of nanocrystalline nickel-tungsten coatings(electrodeposited at 75 °C and 85 °C) at two constant load levels of 5 and10 N with a constant sliding speed of 0.5 m/s. Coatings were applied onan AISI 1045 carbon steel ball. A disk made of AISI 5045 steel with ahardness of 60 HRC was used as the counter face. They reported thatwear rate increased with increasing the applied load from 5 to 10 N inboth coatings. Reduction in grain size up to 15 to 20 nm was accompa-nied initially by a decrease and then increase in wear rate. The wearmechanism for most coatings was adhesive wear while brittle fracturewith sharp microcracks occurred in the hardest coating (640 HV0.1).Haseeb et al. [217] investigated wear and friction of electrodepositednanocrystalline nickel-tungsten coatings with average grain sizes of 20to 22 nm with respect to a ball made of hardened carbon chromiumsteel (100 Cr6). They compared the acquired results with other findingsfrom wear tests on nickel coatings obtained from a sulfonate bath. Ac-cording to their work, a massive transfer of iron was observed overthe nickel surfacewhile therewas a small amount of nickel on the coun-terpart. Contrarily, transfer ofmaterials occurred between both sides foriron and Ni-W couple. Prevailing presence of spherical oxide particles,as well as oxide films over the wear track of the nickel sample, com-pared with the nickel-tungsten sample, decreed the undeniable role ofoxidative wear besides adhesive wear in nickel wearmechanism; how-ever, oxidative wear was also somehow obvious in Ni-W sample. It wasreported that COF of nickel/iron pair was higher than that of Ni-W/iron.Analysis of the worn surfaces revealed no sign of brittle fracture; be-sides, an increase in tungsten content of the nanocrystalline coatingwas followed by a decrease in wear track width and wear rate of thecounterpart. In this regard, two important factors in reducing wearcaused damages are: 1) the increase in hardness due to an increase intungsten content and consequently reduction in contact area betweenthe film and counter body and 2) generation of back transfer towardthe steel ball due to presence of tungsten which forms a protectivefilm retarding further transfer. In a similar research, Amadeh et al.[11], studied wear behavior of nickel-tungsten coatings applied on car-bon steels. They reported that microhardness and wear resistance ofcoatings are improvedwith a decrease in duty cycle and current densitywhich lead to the increased tungsten content. The presence of numer-ous scratches along with adhesive tearing and plow lines over thewear track indicates that thewearmechanism is controlled by amixtureof adhesive wear and abrasion.

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Rupert and Schuh studiedwear properties of nanocrystalline nickel-tungsten alloys with grain sizes of 3 to 47 nm. Their results exhibit thatwear behavior of Ni-W coatings with grain sizes above 10 nm has ac-ceptable correlation with Archard equation (Eq. (17)). It should benoted that the Archard equation breakdown occurs at grain sizesbelow 10 nm [166].

V ¼ K � l � PH

ð17Þ

where, K is called the wear coefficient, l is the sliding distance, P is theapplied load, and H is hardness. It was observed that the worn volumein samples with grain sizes below 10 nm is considerably lower thanwhat expected according to Archard equation (Fig. 27).

Besides, hardness measurements over the wear track revealed thathardness was increased after the wearing process. The hardness in-crease in samples with grain sizes above 10 nm was reported to behigher than thosewith grain sizes below10nm. TEManalyses onmicro-structural evolution near the wear surface (Fig. 28) showed that micro-structure of 25 nm-grained samples did not undergo a noticeablealteration before and after the wear test. However, the microstructureof the near-to-surface area in the sample with grain size of 3 nmchanged throughwearing and coarse grains being up to 20 nm in diam-eterwere developed. Hardening afterwear is primarily due to highly lo-calized deformation and then partial deformation caused by frictionalheating. Hardness increase caused by frictional heating was also attrib-uted to grain growth and grain boundary relaxation.

Wear properties of nickel-tungsten coatings (containing 4.4 at%W)with cracked and crack-free structures were evaluated under dry andoil-lubricant conditions. The cracked coating had a lower coefficient offriction (COF) and wear rate than crack-free coating under both dryand oil-lubricant conditions. Under the dry condition, this observationis explained by the accumulation of debris within cracks so that theywould not act as abrasive particles. It was expressed, regarding oil-lubricant condition, that lubricating medium is stored within cracksresulting in a reduction of wear rate [195].

Wear properties of nickel-tungsten coatings with layered structureand tungsten contents of 2.1 at% (sample A), 2.9 at% (sample B),4.9 at% (sample C), and 11.7 at% (sample D) were investigated [10].Wear results of multilayer coatings with AB/AC/AD structure and thick-ness of 30 μ show that the COF (the average of COF achieved in the dis-tance of 300 to 500 m) was in the range of 0.3 to 0.5 and COF increasedbecause of the high tungsten content of samples. Results also indicatethat COF increases with a reduction in a layer thickness of AC specimenfrom5000 to 25 nm; however,wear resistance is doubled. This fact is at-tributed to the increase in hardness of coatings from 3.7 GPa to 6.7 GPa.Wear rate of monolithic and multilayer nickel-tungsten coatings

Fig. 27.Wear volume versus grain size (a), wear volume versus reciprocal hardness (b), as-depoand thewear track and the as-deposited hardness difference (inset of C). In (a) and (b) the dottethat hardening of the wear track is more pronounced at the finest grain sizes [166].

underwent degradation with an increase in tungsten content. Eventhough it was expected that wear rate of the multilayer coatings wasequal to the average value of those of the two constituting layers (for ex-ample, wear rate ofmultilayer AD is equal to the average ofmonolithic Aand D wear rates), the measured wear rate of layered coatings was lessthan the calculated average.

In order to study the synergistic effect of corrosion and wear onnickel-tungsten coatings, Lee et al. [9] carried out corrosion, wear andtribocorrosion tests, separately, onNi andNi-W samples. Tribocorrosiontests were carried out at−200(VSCE), 400(VSCE) and 600 (VSCE) poten-tials with respect to OCP in a 5 wt% solution. Active, passive, and trans-passive regions were observed for Ni samples while the trans-passivephenomenon was not detected in Ni-W samples. Corrosion productsanalyses determined Ni preferential dissolution and a composition ofNiO, Ni(OH)2, and WO3. The presence of a passive corrosion film onNi-W surfaces not only stabilizes corrosion but also serves a lubricatingagent in tribocorrosion of samples. Moreover, by increasing the poten-tial toward anodic potential (600 mV w.r.t OCP) a reduction in COF isobserved. Tribological studies showed that corrosion and wear interac-tion, particularly at high potentials, affect weight loss in samples signif-icantly. In addition, considering the synergistic effect of tribocorrosion,one may come to the conclusion that effect of corrosion on wear ismore intense than that of wear on corrosion in nickel-tungsten samples.

He et al. [218] added tungsten to Ni-P coatings in order to achievedesired wear properties. They investigated as-deposited and heat treat-ed Ni-W-P coating by means of ball-on-disk apparatus immersed in a3.5% NaCl solution and deionized water. Although the effect of tungstenincrease on tribocorrosion properties was not adequately addressed,their work revealed that abrasive wear was the dominant mechanismin as-deposited and heat treated coatings in deionized water. Besides,it was found that as-deposited coatings undergo adhesivewear and cor-rosion damage in saline media. Meanwhile, heat treated coatings firstshowed abrasive wear and subsequently displayed corrosion damages.A synergistic effect between wear and corrosion in the saline solutionwas observed while it was less noticeable in heat-treated samples.Wear properties of laser cladded Ni-W-Si coatings applied on 45 steelsubstrateswere also determined elsewhere [219]. The coatingswere re-ported to enhance the wear resistance of steel substrates by nine times.COF fluctuation with time was attributed to the presence of extremelyhard W and WSi2 phases as well as the non-uniform structure of coat-ings. In addition, COF and wear rate slightly decreased with an increasein tungsten content leading to less wear resistance.

11.1. Micro/nanocomposite

Incorporation of micro/nanoparticles in the alloy matrix and fabrica-tion of metal matrix composites is another mechanism for enhancing

sited specimens and thewear tracks hardnessmeasurements as a function of grain size (c),d blue linemarks the expectedwear volume based on the Archard equation and (c) shows

Fig. 28. SEMmicrographs related to the lamellae cut from Ni-W as-deposited specimen (a) and the wear track center (b). Bright field TEM images of the Ni-W as-deposited (c) and weartrack lamellae (d) from the specimenwith 25 nm initial grain size; the dashedwhite linemarks the surface of the specimens. (e) and (f), respectively, related to the cross-sectional brightfield TEM of the as-deposited and wear track lamellae from the alloy with 3 nm initial grain size [166].

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tribological properties of nickel-tungsten coatings. For this purpose, thereexist numerous types of particles especially ceramicmicro/nanoparticles.For instance, Yao et al. [220] studied wear behavior of Ni-W-SiC coatings.Theweight loss caused bywear was reducedwith an increase in SiC con-tent in Ni-W coatings. SiC nanoparticles improved strength and hardnessof coatings through dispersion and particle strengthening. By changingthewearmechanism into abrasivewear, a decrease inwear rate and, con-sequently, enhanced wear resistance is encountered. Hence, the abrasivegrooves observed in thewear trackwere attributed to the presence of SiCnanoparticles. In another work, Yao et al. [221] investigated the effect ofalumina particles onwear behavior of nickel-tungsten coatings. Hardnessand wear behavior of nickel-tungsten coatings are severely affected bythe content of alumina nanoparticles. Accordingly, alumina is less impor-tant in wear behavior of coatings when its content is below 3.3 wt%. Onthe contrary, coatings with N3.3 wt% alumina had considerably more fa-vorable wear properties. The presence of ceramic particles such as alumi-na may improve the hardness of Ni-W coatings. For instance, an increasein the content of alumina microparticles from zero to 9.5 wt% led to theenhanced hardness of nickel-tungsten coatings from about 506 Hv toabout 670 Hv. Alumina particles reduce wear rate by forming abrasivegrooves through hwear track. It was reported that alumina nanoparticlesmight act as a solid lubricant and improve wear resistance. In the case ofalumina nanoparticles, it was stated that wear resistance of nickel-tungsten coatings was increased by two to four times with the additionof alumina nanoparticles [222].Microhardness and COFmay be increasedand decreased, respectively, through the distribution of diamond parti-cles in nickel-tungsten alloy matrix. Thus, the presence of these particlesimproves tribological properties of nickel-tungsten coatings [101]. Syner-gistic effect investigation of nanodiamonds, CNT, and alumina particleson wear properties of nickel-tungsten coatings shows that wear resis-tance was improved by an increase in contents of those nanoparticlesthrough the alloy matrix. The following Eq. (18) holds for wear rate andnanoparticles contents:

WR ¼ −1:8048DONð Þ þ 30:858 ð18Þ

where, DON is density of nanoparticles in (#100 × 100 μm2) andWR iswear rate (mm3/N m × 10−5) [223]. The presence of MWCNTs

improved wear resistance by decreasing COF and altering the wearmechanism to abrasivewear [135]. Results of studies on nanocompositeNi–17.5 at%W coatings carried out by Hou et al. [99] reveal that hard-ness of Ni-W coatings was increased from 840 to 860 Hv through an in-crease in the content of alumina nanoparticles from 0.36 wt% to1.59 wt%. The increase in alumina content reduced wear rate by 70%and diminished COF from 0.438 to 0.249. As shown in Fig. 29, coatingsbearing more alumina nanoparticles (a) have a relatively constant andnon-fluctuating friction coefficient curve (COF = 0.249). However,COF curve experienced many fluctuations (b and c) with a decrease inalumina content and its value reached 0.6. Fig. 29c shows that COFhad many fluctuations in the first 240 m of wear distance while itreaches a stable value in the last 160 m. This behavior is explained bythe non-uniform distribution of nanoparticles within the Ni-W matrix.The sudden decrease in COF is attributed to the increase in the contentof nanoparticles within the depth of coating which transforms slidingwear into rolling wear. Addition of more alumina nanoparticles to thematrix reduced COF and enhanced tribological properties of Ni-W coat-ings. It was stated about worn surfaces that the wear track was smoothand material transfer occurred at some points. Plastic deformation andplowing traces were observed in coatings containing theminimum alu-mina content. Tribological properties of Ni-W/PTFE coatings were re-cently evaluated. Increase in PTFE content reduced COF from about0.71 to 0.2 and improved wear properties [216]. As the reinforcing par-ticles in the nickel-tungsten matrix, the presence of graphene oxide[129] and boron nitride [175] was also reported to cause COF drop to0.2 and reduce wear resistance of coatings.

It is worth mentioning that, the addition of some nanoparticles maydegrade the tribological behavior of Ni-Wcoatings by the adverse effect.In this regard, studying the effect of MoS2 addition to nickel-tungstenmatrix reveals that MoS2 particles increase roughness and averagegrain size of coating, and led to decreasing microhardness, tungstencontent, and adhesion of the coating to the substrate. Molybdenum sul-fide nanoparticles caused irregular tribological properties in nickel-tungsten coatings. It was reported that traces of MoS2 in nickel-tungsten coatings halved friction coefficient from 0.27 to 0.14. The in-crease in MoS2 content led to an irregular wear behavior as well as a re-duction in hardness and adhesion of coatings. High MoS2 contents

Fig. 29. FOC versus sliding distance for Ni–40wt%W-Al2O3 nanocomposite coatings with1.59 wt% (a), 0.7 wt% (b), and 0.36 wt% (c) Alumina nanoparticles [99].

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resulted in a sponge structure and abnormal changes in COF which in-creased up to 0.47 with coatings removal [224]. Similar behavior wasalso reported for zirconia addition. Firstly, wear behavior was improvedby the addition of ZrO2 nanoparticles which increase hardness andstrength of coatings caused by dispersion and particle strengtheningmechanisms [103]. Although wear resistance increased with zirconiacontent rising up to 5wt%, it was reported that further increase in zirco-nia content by 12wt% diminishes thewear resistance as a result of coat-ing embrittlement [70].

Regarding the tribological properties of Ni-W coatings at high tem-peratures, lanthanum oxide microparticles brought about considerablechanges in wear behavior of nickel-tungsten alloys at 700 °C. Sincelanthanum oxide is employed as filler and reducer of COF inpolytetrafluoroethylene (PTFE) composites, it may enhance tribologicalproperties of coatings at high temperatures. COF of coatings (with re-spect to glass) was reduced to the minimum value of 0.22 with an in-crease in La2O3 content up to 7 wt%. The improvement in tribological

properties caused by the addition of lanthanum oxide was reported tobe due to the hexagonal structure of La2O3 layers and their chemical sta-bility at 700 °C [205]. The addition of CeF3 nanoparticles was consideredan appropriate approach to enhance wear behavior of nickel-tungstencoatings at high temperatures. Assessing the wear resistance of coatingsat 650 °C against molten glass showed that fluctuations of COF is fadedand its value is dropped from 0.43 to 0.18 with an increase in nanoparti-cles content from zero to about 5–6 wt%. This behavior is explained bythe lamellar-hexagonal structure of CeF3 nanoparticles as well as theirchemical stability at high temperatures [172].

Investigation of high temperature (400–600 °C) wear behavior ofNi–W–Co/SiC composite brush plated coatings using a plate-on-ringtest rig shows that at 500 °C, the composite coating exhibits the lowestwear and COF in the test. Comparatively, at high temperatures, bothabrasive and adhesive wears are the major wear mechanisms [225].Hu et al. [226] studied the performance of high-temperature (500–600 °C) TiN coatingswith a supporting and adhesive nickel-tungsten in-terlayer. It was reported that TiN coatings applied on a nickel-tungstenlayer had a lower COF and wear rate at high temperature; with a Ni-Wlayer, the dominant wear mechanism was a combination of adhesiveand abrasive wear. Tables 3 and 4 provide a brief representation ofwear mechanism and wear data related to the different Ni-W samplesand conditions.

12. Heat treatment

12.1. Microstructural evolution

Heat treatment of nickel-tungsten alloys has been performed in arelatively wide temperature range (100–1000 °C). Temperature, an-nealing duration time, and alloy chemical composition (according tothe nickel-tungsten phase diagram) have important roles in heat treat-ment and microstructure of coatings. Although the low-temperatureannealing (below 250 °C) was reported to have almost no effect onamorphous structure of Ni–17–19 at%W alloys, some other investiga-tions revealed that annealing for 1 h even at 100 °C might somehowchange the structure of nickel-tungsten alloys [127,145]. Itwas reportedthat heat treatment at 400 °C for 1 h did not cause any significant chang-es in phase structure of Ni-22.1 at%W; except grain size increased from4.8 to 5.9 nm. There are similar studies that show heat treatment up toabout 400 did not affect considerably Ni-W structures [19,55,101,228].It is reported that the increase in grain size was intensified with an in-crease in annealing temperature from 400 to 500 °C [115]. Crystalliza-tion occurred after 500 °C and the peak intensity of Ni (111) planewas higher than as-deposited samples and the Ni4W peak appeared insamples annealed at 500 °C. The crystalline structure of Ni4W phase(Fig. 30) consists of superimposed body-centered tetragonal W (a:5.73 Å, c:3.55 Å) and face-centered tetragonal Ni (a: 3.62 Å, c:3.55 Å)structures [161].

Heat treatment of nanocrystalline Ni-W alloys at 600 °C or above ledto precipitation of Ni4WandNiW intermetallic phases dependingon thechemical composition (according to the phase diagram). However,structure evolution and hardness of nickel-tungsten alloys are stronglyaffected by annealing temperature and chemical composition (phase di-agram). Considering the chemical composition based on the nickel-tungsten phase diagram, the nickel alloy with 7 at% tungsten is locatedin the region of tungsten solid solution in the nickel where an increasein annealing temperature and time leads to the formation of coarsegrains. Besides, the nickel alloy with 14 at% tungsten is located in the(Ni) and Ni4W two-phase region where an increase in annealing tem-perature leads to the growth of Ni4W intermetallic precipitates in the(Ni) matrix and consequently increase in hardness of alloys [127,227,229,230]. A series of XRD patterns related to the heat treatment of Ni–21 at%Wat temperatures of 300, 450, 600, 750, and 900 °C are presentedin Fig. 31.

Table 3Wear mechanism, in brief, in different Ni-W samples.

Samples Test Reason Mechanism Ref.

As deposited Wear Fluctuations in COF, ferrous oxide and plastic deformation in worn surfaces Adhesive [64]Annealed at 300, 500, 700 Wear Less fluctuation in COF, smooth worn surface, no plastic deformation. Abrasive [64]As deposited Wear Traces of wear groove due to abrasion of debris Abrasive [8]Annealed at 700, 900, 1100 Wear Transfer layer, cracks formed in the wear tracks Adhesive [8]As deposited Wear Surface plowing, and transfer layers, and cracks Adhesive and abrasive [10]Multilayer Wear Scratching marks and transfer layers Adhesive and abrasive [10]As-deposited (in saline) Tribo-corrosion Oxide layer, small amount of peeling, pits and shallow scratches (severe) Adhesive and corrosion [218]Annealed at 300, 400, 500, 600 (in saline) Tribo-corrosion Oxide layer, small amount of peeling, pits and shallow scratches Abrasion and corrosion [218]

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Microstructural evolution studies on amorphous/nanocrystallineNi–20–21 at%W alloys revealed that these alloys have a very good ther-mal stability up to 500 °C. Nevertheless, the thermal stability of theamorphous phase was dependent on chemical composition and thespecific features of a microstructural network of alloys [145]. Detorand Schuh [231] investigated grain size evolutions, segregation, andgrain boundary structures in Ni-W alloys in a wide range of grain sizes(3 to 70 nm). The grain size of nanocrystalline Ni-W alloys did notchange with heat treatment at 450 °C for 24 h with the considerablethermal stability of the structure. However, normal grain growth oc-curred at higher temperatures. It is clear that thermal stability of nano-crystalline Ni-W is a function of heat treatment temperature and timeand initial grain size (Fig. 32). Therefore, the following sequence maybe expected for heat treatment of nanocrystalline Ni-W alloys: grainboundary relaxation, grain growth, grain boundary segregation, andchemical ordering.

For instance, Ni-W alloy structure with an initial grain size of 3 nmremained almost unchanged at 450 °C with an excellent thermal stabil-ity. Although no sign of grain growth was observed at lower annealingtemperatures, an increase in structural hardness was observed. The re-leased heat in differential scanning calorimetry at low temperatures(below 300 °C) was attributed to grain boundary relaxation. The con-served energy is released during relaxation in the form of excess grainboundary dislocations as the boundary transforms to amore stable con-figuration. Since grain boundary structure plays a significant role innanocrystalline materials, grain boundary relaxation leads to enhancedstructural strength. Hence, an increase inmicrohardness of nanocrystal-line nickel-tungsten alloys caused by heat treatment at low tempera-tures is attributed to grain boundary relaxation. Studies on the effectof heat treatment on dynamic mechanical properties of Ni–15.2 at%Walloy indicated that the best dynamic strength and ductility wasachieved by annealing at 300° [115]. Temperature rise up to 600 °Cand grain-coarsening, depending on the chemical composition, lead toprecipitation of secondary phases (like Ni4W) in the alloy matrix. TEM

Table 4A brief wear data related to the different Ni-W samples and conditions.

Specimen Counter body Wear type Load Speed

Ni-W WC sphere Pin on disk 5 N 0.15 mNi-W Hardened steel Pin on disk 1 N 0.03 mNi-W Steel hardened to 63 HRC Pin on disk 5&10 N 0.5 mNi-W 52100 tool steel Pin on disk 1 N 0.75 mNi-W JIS SKD-11 tool steel 62HRC Ring on disk – –Ni-W wC sphere Ball-on-disk 5 N 200 rNi-W WC ball Ball-on-disk 5 N 200 rNi-W Alumina cube Block on ring 29.4 N 0.21 mNi-W Steel St52 Pin on disk 20 N 0.3 mNi-W GCr15 stainless steel Pin on disk 20 N 5 HzNi-W Uncoated stainless steel Pin on disk 6 N 1.03 mNi-W Glass High temp. tribometer 10 N 1.5 mNi-WNi-W-P

Cr bearing steel (GCr15) Ball-on-disk 150 N 150 r

Ni-W/SiC Cr Abrasion test 50 N 400 rNi-W/Al2O3 – Abrasion test 50 N 401 rNi-W/diamond Silicon nitride Si3N4 Ball-on-disk 10 N 300 rNi-W/La2O3 Glass High temp. tribometer 10 N 1.5 m

micrographs and XRD spectra of samples heat treated at 600 °C orabove revealed that grains grew with an increase in annealing durationor annealing temperature; though the grain growth rate in respect toannealing temperature strongly depends on the initial grain size. Infact, grain growth rate is extremely low in fine-grained alloys (Fig. 33).

After grain growth, grain boundary segregation and chemical order-ing transitions occurred in nanocrystalline nickel-tungsten alloys withan increase in annealing temperature. Atomprobe tomographyanalysesshowed that tungsten was not uniformly distributed in coating struc-ture of samples (20 at%W) heat treated at 600 °C and its segregationwas vividly observed. According to the measurements, the averagetungsten content in grain boundaries and interior areas was about 19and 25 at%, respectively [232,233]. Fig. 34 shows TEM micrographs ofnanocrystalline Ni–20 at% W alloys before and after heat treatment aswell as atomicmap schema of tungsten atoms position. Chemical order-ing transitionwas also observed after grain growth in samples annealedat sufficiently high temperatures. Chemical ordering generally beginswith the disorder in the as-deposited state to short range orderingand subsequently follows long range ordering with precipitation of in-termetallics [231].

Although it was reported that grain size increased at high heat treat-ment temperatures, comparing nickel-tungsten and nickel-tungsten-diamond alloys revealed that grain size of Ni-W-diamond samples in-creased less than Ni-W samples after heat treatment [101].

12.2. Heat treatment effects on mechanical properties

Annealing effects on mechanical properties of Ni-W alloys dependon the alloy composition, heat treatment temperature, and time. Inthis regard, there are various studies in the literature which investigatesdifferent annealing temperature on Ni-W alloys with several alloy com-position. For instance, microhardness of Ni–7 at%W alloy was observedto continuously fall with an increase in annealing temperature from 100to 600 °C (Fig. 35). Contrarily, this increase in temperature enhanced

Temp Humidity COF Hardness Type Ref.

/s r.t. – 0.60–0.67 4.0–7.1 GPa Tribology [166]/s r.t. 50% 0.73–0.82 535–567 HV0.05 Tribology [217]

/s r.t. 60–70% 0.4–1.2 575–640 HV0.1 Tribology [67]/s r.t. – – 514–537 HV0.08 Tribology [11]

r.t. – 0.35–0.45 594 to 823 Hv0.1 Tribology [64]pm – – 0.50–0.52 7.0–10.5 GPa Tribology [8]pm – – 0.3–0.5 3.7–6.7 GPa Tribology [10]/s – – 0.19–0.31 630–660 Hv0.1 Tribo-corrosion [9]

/s – – – 693–1258 Hv Tribology [143]5 mm r.t. – 0.50–0.62 878–946 HV0.05 Tribology [195]/s r.t. – 0.2–0.3 650 HV0.05 Tribology [224]

m/s 700 – 0.37–0.45 587 HV0.2 Tribology [205]pm r.t. – 0.3–0.6 552–860 HV0.2 Tribo-corrosion [218]

pm r.t. – – 506–711 HV0.05 Tribology [220]pm r.t. – – 506–689 HV0.05 Tribology [221]pm – – 0.55–0.65 710–805 HV0.1 Tribology [101]m/s 700 – 0.25–0.3 650–752 HV0.2 Tribology [205]

Fig. 30. Crystal cell of Ni4W intermetallic phase [161].

Fig. 32. Heat treatment map for nanocrystalline Ni-W and pure electrodeposited Ni. Thismap depicts the expected sequence of grain boundary relaxation, grain growth, grainboundary segregation and chemical ordering [231].

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microhardness of Ni–14 at%W alloys [127]. Fig. 36 compares variationsin Vickers hardness in brittle and soft Ni-W alloys as well as pure nickelsamples all annealed at different temperatures. The same behavior wasalso reported for the effect of annealing temperature on the hardness ofNi-W-Fe coatings in which the best mechanical properties were ob-served in samples annealed at 500 °C [162,234]. Isochronal heat treat-ment of Ni-W alloys at average annealing temperatures and under avacuum generally leads to an increase in hardness of tungsten solid so-lution in FCC-Ni. For instance, it was reported that heat treatment at400 °C for 1 h did not cause any significant changes in phase structureof Ni-22.1 at%W; whereas grain size increased from 4.8 to 5.9 nm andhardness rose from 579 to 594 kg mm−2 [19,228]. Auerswald andFecht measured the microhardness of as-deposited Ni-W coatings else-where at 693 ± 10. It was stated that Vickers hardness increased up to1258 ± 19 HV with an increase in annealing temperature by 600 °C;while further increase in annealing temperature up to 800 °C led to a de-crease in Vickers hardness falling to 800 ± 19 HV. Hardness increasewas assigned to increase in ordered nanocrystalline phase (grains) tointercrystalline phase (boundary layer phase) ratio. On the otherhand, hardness decrease in samples annealed above 600 °C was attrib-uted to the considerable grain growth in coatings. The grain size of theelectroplated nickel-tungsten alloy was 6 nm which increased to35 nm with an increase in annealing temperature up to 800 °C [143].It was reported that nitrogenation process at 400 and 500 °C increasedhardness and, in some cases, inter-atomic distance in solid solution lat-tice of nickel-tungsten alloys [235].

Fig. 31. XRD spectra related to the heat treatment of Ni–21 at%W at temperatures of 300,450, 600, 750, and 900 °C [231].

The effects of heat treatment on hardness and wear resistance of Ni-W coatings have been frequently addressed in literature and it was re-ported that nickel-tungsten alloysmaintained acceptablywear resistantup to 500 °C [38]. Hou et al. [64] heat treated nickel-tungsten sampleswith a heating rate of 10 °C/min at 300, 500 and 700 °C for 1 h undervacuum.Microhardnesswas improved by an increase in annealing tem-perature from 300 to 500 °C; while further increase in annealing tem-perature up to 700 °C led deterioration in coatings microhardness.This behavior is described with grain coarsening in the microstructure.Adhesive wear mechanismwas reported to be prevailing due to fluctu-ations in COF (0.35 to 0.45) and the presence of plastic deformation andferrous oxide over the worn surface of as-deposited specimens. Howev-er, as precipitation of intermetallic compounds such as Ni4W and NiWmay help the hardness improvement at temperatures above 300 °C, adifferent wear mechanism was proposed for coatings heat treated at700 °C due to low fluctuations in COF (below 0.3) and no sign of plasticdeformationover theworn surface. Experimental results indicatedwearresistance enhancement because of the increased hardness according toArchard's law. Coatings containing 8.3 at%Wand grain size of 3 nmweresubject to an annealing process elsewhere at 700, 900, and 1100 °C for0.5 to 3 h under nitrogen atmosphere. The hardness of as-depositedcoatings was about 7 GPa. It decreased from 10 GPa to 7.5 GPa with anincrease in annealing temperature from 700 to 1100 °C. Grain size alsoincreased from 44 to 70 nm within the same temperature range. Be-sides, wear resistance of heat treated Ni-W samples was lowered byan increase in annealing temperature from 700 to 1100 °C. However,

Fig. 33.Measured grain size using XRD analysis after 24-h heat treatments related to threenanocrystalline Ni-W alloys with different initial grain sizes (at 25 °C), results showexcellent stability up to ∼500 °C [231].

Fig. 34. TEM images of as-deposited (a) andheat treated for 3 h at 600 (b) nanocrystallineNi-Walloy. (c) Exhibits an atommapplotting the position ofWatoms in a 4-nm-thick slice of theheat treated specimen [232].

1002 M.H. Allahyarzadeh et al. / Surface & Coatings Technology 307 (2016) 978–1010

the hardness of annealed samples was still higher than that of as-deposited samples, implying a breakdownof Archard relation. The abra-sive wearmechanismwas assigned for as-deposited samples due to thepresence of wear grooves over the worn surfaces caused by abrasion ofdebris. However, the presence of transfer layers and initiated cracks im-plied the adhesive wear mechanism for annealed coatings. Microhard-ness measurements showed that the hardness over the wear track inas-deposited samples increased after the wear test due to the grain re-laxation induced by wear heat. In comparison, the microhardness in-crease in annealed coatings was insignificant [8].

He et al. [218] showed that, in the case of Ni-W-P coatings, hardnessinitially increased from 635 to 875 Hv with an increase in heat treat-ment temperature up to 400 °C and subsequently dropped to 645 Hvwith a further rise in heat treatment temperature from 400 to 600 °C.According to wear behavior analysis of Ni-W-P coatings, the minimumwear ratewas observed in coatings heat treated at 400 °C. It was report-ed that heat treated samples had the least wear rates and the synergisticeffect of corrosion and wear in heat treated samples in a saline solutionwas less noticeable than that in as-deposited samples. Analyses on mi-crohardness and corrosion resistance of electroless Ni-W-Cr-P alloycoatings showed that the highest microhardness and corrosion resis-tance was achieved in samples annealed at 800 °C [236].

Wang et al. [101] reported that hardness values of Ni-W and Ni-W-diamond samples increased by 1160 and 1205 Hv, respectively, with anincrease in annealing temperature. This increase might be because ofNi4W precipitation. However, hardness diminished with the increasein annealing temperature up to 700 °C due to grain coarsening. Weartests indicated that, according to Archard relation, the wear rate was

Fig. 35. Microhardness of Ni–7 at%W and Ni–14 at%W coatings annealed at differenttemperatures [127].

inversely related to the hardness of coatings. The presence of diamondin Ni-Wmatrix led to the decreasedwear rate and COF and consequent-ly an improvement in wear resistance. Moreover, considering theweight loss results for Ni-W and Ni-W-diamond samples, wear rate de-creased with an increase in annealing temperature up to 600 °Cwhile itunderwent deterioration with further increase in annealing tempera-ture up to 700 °C. SEM micrographs of the worn surfaces illustratedplastic deformation, abrasive grooves, plowings, and scratches whereasthe depth of wear grooves in composite samples was reported to be ex-tremely low. The aforementioned wear signs observed in as-depositedsamples were not found on the worn surface of heat treated samplesat 600 °C. In addition, the worn surfaces of diamond-composited sam-ples were relatively smooth. It was reported that heat treatment im-proved the hardness of Ni–1.35Re–5.46W–3.82B–15 wt%SiC coatingsfrom 600–800 Hv to 1400–1650 Hv. The highest wear resistance ofthis coating was acquired by heat treating at 500 °C which was 4.5times more than that of hard chromium. The synergistic effect of heattreatment temperature and chemical composition on hardness andwear resistance of coatings is illustrated in Fig. 37a and b [237].

12.3. Thermal stability

Thermal stability of nickel-tungsten alloys is strongly dependent onchemical composition andmicrostructure so that Tx (crystallization tem-perature) increased at high tungsten contents. Generally, nanocrystallineNi-W alloys demonstrate excellent thermal stability up to 500 °C [231].Thermal stability of two groups of nickel-tungsten alloys (17–19 at%W,

Fig. 36. Comparison between Vickers microhardness in soft and brittle Ni-W alloys andelectrodeposited pure Ni at various anneal temperatures [229].

Fig. 37. Effect of heat treatment on hardness (a) and wear resistance (b) of Ni–W–2.75%B–12%SiC (♦), Ni–W–3.8%B–12.2%SiC (■), and RE–Ni–W–3.8%B–15.2%SiC (▲) [237].

1003M.H. Allahyarzadeh et al. / Surface & Coatings Technology 307 (2016) 978–1010

type a and 21–23 at%W, type b) was investigated by Grabchikov et al.[145]. They observed that thermal stability of type-a films with a homo-geneous structure was 150 to 200 °C higher than type-b films with acoarse-grained network structure. Through studying thermal stability inNi-Walloys containing 5 and15 at% tungstenHajTaieb et al. [35] reportedthat Ni-5 at%W alloy maintained microhardness at 400 °C after 16 h;however, it dropped severely at 700 °C after a short moment and de-creased by 50%. In the case of Ni-15 at%W,microhardness dropped insig-nificantly at 700 °C after 16 h. Regarding Ni-21 at%W coatings, thermalstabilitywas reported to be themaximum till 500 °C and recrystallizationalong with the formation of stable Ni4W and NiW phases commenced ataround 600 °C [229]. Another reason for the thermal stability of Ni-W al-loys at high temperature ismight be the formationof oxide particle impu-rities and W-segregations at oxide-Ni(W) interfaces. Nanoscale W-richoxide particles identified in annealed (at 700 °C, Ar-5%H2) Ni-23 at%Wlimited the grain growth. Also, high-resolution TEM analysis shows thatW-segregation occurs at the oxide/Ni(W) interfaces and it is not ob-served only at the Ni(W)/Ni(W) grain boundaries as predicted by ther-modynamic models. Experimental results show that impurity oxideparticles hindered grain growth and improve nanoindentation hardness;hence, impurity phases should not always be deemed as detrimental tonanocrystalline thermal stability [238]. The critical influence of carbonon the thermal stability of nanocrystalline Ni–23 at%Walloyswas also in-vestigated. Since carbon has a low solubility in theNi(W) solid solution, itmay segregate strongly at the grain boundaries whileW anti-segregationmay occur during carbide precipitation. In this regard, carbon contamina-tion may affect the thermal stability of nanocrystalline Ni–W [239].

The addition of various elements to the Ni-W binary system couldhave an efficient effect on the thermal stability of these alloys. For exam-ple, experimental analyses on Ni-Co-W alloys revealed that the presenceof cobalt led to a reduction in thermal stability of Ni-W films [145].

13. Modern trends in fabrication of novel Ni-W films

Recently, a large number of studies have been published regardingtungsten electroplating with iron group elements [240]. This section,however, sets out to specifically address the novelmeasures undertakenin the field of Ni-W alloys.

13.1. Diffusion barrier

In order to produce diffusion barrier coatings and restrict inter-diffusion between the superalloy and aluminide coating, Cavalletiet al. [241,242] applied a 7 μm-thick nickel-tungsten layer over 4th-generation nickel-based superalloys. The film was annealed at 1100 °Cfor 16 h prior to aluminizing and formation of an aluminide diffusioncoatings. It was observed, after aluminizing, that a film consisting W-rich precipitates was formed between the superalloy and aluminidecoating (and platinum-aluminide) which modified oxidation behaviorof the diffusion coating and impeded inter-diffusion. It was reportedthat the diffusion barrier film reduced β-NiAl → γ′-Ni3Al phase

transformation and inhibited the formation of secondary reactionzones (SRZ) in the superalloy/coating interface. In a similar researchwork, Re-(Ni-W) compounds were employed to produce a diffusionbarrier film and restrain inter-diffusion from the substrate alloying ele-ments toward the coating. Herein, a 2 μm-thick Re film and then a 7 μm-thick Ni-24.2 at%W film was applied over the superalloy. EDS line anal-yses after oxidation show that this coating has a desirable performanceas a barrier against aluminum diffusion from the coating toward thesubstrate and also alloying elements diffusion from the substrate to-ward the coating [243]. In this work, Ni-24.2 at%W coatings werenamed as diffusion barrier films. According to Ni-W phase diagram,this amount of tungsten corresponds to the two-phase region consistingNi4W and NiW. Heat treatment before application of aluminide/platinum-aluminide coatings may facilitate the formation of the afore-mentioned intermetallic phases in the diffusion barrier film. Restrictionof inter-diffusion of alloying elements from the substrate toward thecoatings and vice versamay be due to the formation of such intermetal-lic phases. Regardingdiffusion inNi-W solid solution, there is some pub-lished literature for further information [244–246].

13.2. Ni-W alloys and nanocomposites with gradient structures

Recently, many researchers have succeeded in electrodeposition ofNi-W alloy and nanocomposite coatings by means of various tech-niques. Wang et al. [192] produced functionally graded coatings withinsix steps by a gradual and simultaneous increase in temperature andcurrent density. Their results showed that tungsten content rose fromthe substrate, 3.4 at%, to the surface, 28.1 at% (Fig. 38). They showedthat crystallite size and lattice strain decreased and increased, respec-tively, from the substrate to the surface. Schuh et al. [144] producedsuch coatings by means of the pulse reverse technique (Fig. 39a). Theydecreased tungsten content to the surface by increasing the reverse cur-rent density and showed that nanocrystalline FG coatings could be pro-duced in a way that grain size increased toward the surface whilehardness is decreased. The coatings structure was amorphous near tothe substrate and then grain size increased from 6 nm to 60 nm. Ni-W-Al2O3 nanocomposite coatings were recently prepared on steel sub-strates by means of pulse current technique. These coatings were pro-duced applying diffusion-control kinetics of tungstate ions andalumina nanoparticles which existed in the electrodeposition solutionat extremely lower concentrations. Tungsten and alumina contents inFG nanocomposite coatings increased gradually to the surface [110].

13.3. Ni-W alloys and nanocomposites with multilayer structures

Although multi-layered nickel-tungsten structures have been al-ready produced through other processes such as ion beam sputtering[247,248], different techniques are employed for electrodeposition ofNi-W coatings. Nickel-tungsten amorphous/nanocrystalline layeredstructure was prepared by means of reverse current technique (Fig.39b) [10,144]. The reverse current was alternatively changed to

Fig. 38. Cross-sectional backscatter electron SEM micrographs (a) and elemental distribution of alloying elements along coating thickness (b) of FG-Ni-W [192].

1004 M.H. Allahyarzadeh et al. / Surface & Coatings Technology 307 (2016) 978–1010

maximum and minimum values. The increase in reverse current corre-sponds to more dissolution of tungsten and consequently grain-coarsening. On the other hand, the minimum reverse current led tothe formation of W-rich films with fine grains. In another research,thin nanoscaled films [191] and thickmicroscaled films [209] were pre-pared by alternatively changing of the cathodic current between 1 and 4A dm−2. Improvement in mechanical and corrosion properties was re-ported for the thin film and produced the coating, respectively, becauseof the formation of a layered structure. Improvement in corrosion resis-tance of the layered structure is attributed to the increased number ofinterfaces responsible for the lateral spreading of the corrosive speciesbetween layers; which in turn leads to the direct attacking of the sub-strate as in the monolithic coating [209]. Multi-layered amorphous Ni-P-W coatings consisting thin layers of Ni-8wt%P–15wt%W and Ni-5wt%P–45wt%W were previously electroplated by Papachristos et al.[249] through alternatively changing of cathodic current between 2and 20 A dm−2. Microstructural and crystallization behavior of coatingswith an increase in annealing temperature from 200 to 800 °C wasstudied.

Fig. 39. Functionally graded (a) and multilayer (b) Ni-W alloys coatings obtained by pulse revvariations as a function of coating thickness [144].

13.4. Hydrophobic coatings

It was recently reported that incorporation of PTFE particles tonickel-tungsten coatings may improve hydrophobic properties. Thecontact angle of Ni-W coatings was about 98.2° which increased up to109.9° with the addition of self-lubricating polymeric particles (PTFE)[216].

14. Mathematical modeling and computer simulation

Researchers and scientists have made many attempts to modelphysical/chemical behavior and carry out computer simulations in thehope of understanding the dominant phenomenon and properties ofsystems and microstructure, respectively. Regarding electrodepositionof Ni-W and also characteristics of the electrodeposited nanocrystallineNi-W, Podlaha and Landolt [77] developed a steady-state mathematicalmodel to predict the induced codeposition behavior of NiwithMo orW.Through this process, they analyzed the kinetics (simple Tafel expres-sion) and mass-transport controlled regions on rotating cylinder

ere technique. Images beneath the SEM micrographs show microhardness and grain size

1005M.H. Allahyarzadeh et al. / Surface & Coatings Technology 307 (2016) 978–1010

electrodes. They assumed a Nernst boundary layer representation forthe mass transfer of ions through the diffusion layer and soluble nickelacting as a catalyst to the generation of an absorbed intermediate spe-cies on the cathode surface. The governing equations in this model arethe steady-state material balance equations in the diffusion layer forWO4

−2 and NiCit− species:

∇ � NNiCit− ¼ 0 ð19Þ

∇ � NWO2−4

¼ 0 ð20Þ

Constant pH in the diffusion layer, obeying charge-transfer kineticsfrom the Butler-Volmer equation, and Langmuir adsorption type arethe assumptions of this model. By applying the boundary conditionsand substitution of electrochemical relations of partial current densityof Ni, W, and water surface coverage, θ, of adsorbed intermediate canbe derived mathematically as following based on the inducedcodeposition mechanism (discussed in Section 3):

θ ¼kc;W;1C

sWO−2

4CsNiCit−1 exp

−αc;W;1nW;1 FRT

E� �

kc;W;1CsWO−2

4CsNiCit−1 exp

−αc;W;1nW;1 FRT

E� �

þ kc;W;2 exp−αc;W;2nW;2 F

RTE

� �ð21Þ

where, kc ,W ,1 is cathodic rate constant of intermediate formation; kc ,W ,2

is cathodic rate constant of intermediate reduction; CWO4−2

s is concentra-tion of tungstate ions; CNiCit−1

s is concentration of citrate ions; αc ,W ,1 iscathodic transfer coefficient in intermediate formation; −αc ,W ,2 is ca-thodic transfer coefficient in intermediate reduction; nW ,1 is numberof transferred electron in intermediate formation; nW ,2 is number oftransferred electron in intermediate reduction; F is Faraday constant;E is potential; R is universal gas constant; and T is absolute temperature.Since at each given potential it is possible to calculate the partial currentdensity of each species, we are also able to calculatemathematically thecathodic current efficiency and chemical composition of deposited alloy[77,240].

Kimoto et al. [250] modeled the deposition rate of Ni-W by the fol-lowing equation at an arbitrary interval Δt and stoichiometrically esti-mated the average current efficiency ē:

Δwα ¼ I 1−αð ÞMNi þ αMW½ �Δt2 1þ 2αð ÞF e ð22Þ

where, Δwα is the deposition rate, α is the W atomic ratio in the deposit(0bαb 1), andMNi andMWare the atomicweight ofNi andW, respective-ly. Franz et al. [130] stated that in pulse electrodeposition the rate ofWco-deposition is governed by the deposition rate of Ni. They assumed that atpulse-on times exceeding the transition time of Ni (where the surfaceconcentration of Ni ion becomes essentially zero) the deposition efficien-cy is expected to fall because of increasing hydrogen evolution reaction.They estimated the dimensionless transition time τ⁎ (τ⁎=Dτ/δ2) as afunction of applied dimensionless current density ip⁎ ( ip

� ¼ ip=iL ) forτ⁎ N 0.1 and at small duty cycles as follows:

τ� ¼ 4π2 ln

π2

81−

1i�p

!" #ð23Þ

where, D is the diffusion coefficient, δ the steady state diffusion layerthickness, τ the transition time, and iL the steady-state limiting currentdensity. This estimation suggests that the deposition rate of tungstenwas affected by the rate of transport of nickel which is the highest atshort times.

In the recent decade, mechanical properties, grain boundary segre-gation, chemical ordering, thermal stability, and dislocation interactionsin nanocrystalline Ni-W have been thoroughly investigated using

computer simulation. In these studies, molecular dynamic simulationand Monte Carlo simulation have been utilized for better understatingthe aforementioned characteristics. Lund and Schuh [251] carried outsimulations of the multiaxial deformation of very fine nanocrystalline(d = 2–4 nm) under various loading conditions. Their results showthat for all studied grain sizes amarked increase in strength under com-pressive loading as compared to tensile loading. They conclude that themagnitude of the strength asymmetry was generally seen to increasewith increasing grain size, suggesting a peak in the compression–tension strength asymmetry at finite nanocrystalline grain sizes. In an-other research, Deter and Schuh [252,253] grain investigated boundarysegregation, chemical ordering, and stability of nanocrystalline nickel-tungsten alloy using atomistic computer simulations. Employingmolec-ular dynamic and Monte-Carlo simulations their findings showed thatthe extent of tungsten segregation in inter-crystalline regions is mostpronounced in thedilute limit and decreaseswith increasing solute con-tent. Simulations also exhibited that short-range chemical ordering isevident in all nanocrystalline structures, and chemical ordering be-comes more pronounced by an increase in the grain boundary density.Calculations also illustrated that grain boundary energy decreaseswith the addition of W, but does not reach zero for any compositionor grain size considered here. Meraj et al. [254] investigated the role ofW on the dislocation evolution in Ni-W alloy during tension followedby compression loading using molecular dynamic simulation. Theirfinding showed that addition of W in Ni lattice decreases the crystallin-ity and results in degrading the deformation dependence on dislocationcontrolledmechanisms. The simulation results at a strain rate of 108 s−1

and temperature of 300 K exhibited that deformation of Ni–15 at%Wduring compressive loading is significantly controlled by dislocationmovement.

Cavaliere [255,256] studied the mechanical properties of plasticallygraded nanocrystalline Ni-W (at 20 and 100 nm constant grain size)via multi-step nanoindentation and finite element calculations. The fi-nite element investigation carried out usingABAQUS simulations exhib-ited that the grading the microstructure and plastic mechanicalproperties of these alloys leads to several changes in the driving forcefor fracture represented by the different behaviors of the effective J-integral at the crack tip.

15. Conclusions and future research directions

We prepared this review paper on Ni-W films to outline parameterincluding electrodeposition conditions, effective characteristics param-eters, corrosion, and high-temperature oxidation, mechanical proper-ties, magnetic features, and thermal stability of these nanocrystallinealloys. We point up the broad trends in the worthy application of Ni-W in different industries. Nickel-tungsten alloys are introduced as aproper and sound replacement candidate for hard chromium deposits.The unique, special, and premium characteristics of these nanocrystal-linefilmsmake them as the cost-effective first choice inwear and corro-sion resistant coatings, electrocatalytic films for hydrogen energy andelectrolysis, micromechanical and microelectrochemical devices, high-temperature stable films, etc. In this paper, attempts were made to col-lect and present a concise information regarding effective parameters,different utilized electrolytes, deposition mechanism, various platingtechniques, and different characteristics of Ni-W nanocrystalline alloys,and micro/nanocomposite deposits.

• Pulse reverse deposition. As discussed in Section 5.3 introduction of ananodic reverse short pulse enables one to electrodeposit fine andsound nanocrystalline nickel-tungsten films through removing andsuppressing hydrogen evolution on the cathode surface. However,this technique is already known and employed in some research;the lack of extension of fundamental knowledge in deposition mech-anism and specified procedure to obtain an optimized film is clearlyevident.

1006 M.H. Allahyarzadeh et al. / Surface & Coatings Technology 307 (2016) 978–1010

• Electrocatalytic properties. Since Ni-W alloys could suppress highertemperatures, it seems that considering its good electrocatalytic prop-erties, this alloy could be utilized at some processes which need ahigher temperature. Incorporation of some catalytic nanoparticlesand employment of specific deposition condition, which lead to typi-calmorphologywith the higher effective surface, can be considered asfuture directions in this trajectory.

• Nanocomposite coatings. Incorporation of nanoparticles, as discussedin previous sections, would lead to obtaining films with specific char-acteristics such as high hardness, premiumwear resistance, good cor-rosion resistance, appropriate mechanical properties, suitablemicrostructure, and acceptable thermal stability. Use of horizontalcathode configuration allows obtaining Ni-W deposits comprising64 wt% embedded diamond nanoparticles (with 2249 ± 23 Hv)[177]. Thus, developing methods for incorporating a higher amountof nanoparticles such as B4C and etc. to achieve very hard and/or cor-rosion/oxidation resistant coating can be supposed as the futuretrends of these investigations.

• Tailored and patterned structures. Tailoring and patterning novel struc-tures using intelligent and targeted deposition conditions can beassigned as future directions. The novel structures such as gradientor layered patterns could be nominated for specific applications.About the multilayer structure, investigation of layer thinning up tonanoscale on mechanical properties is suggested. In this concern, itis possible to develop nanocomposites like structures with thick,nickel-rich solid solution, and soft matrix and very thin W-rich tenu-ous. Heat treatment of such structures which lead to precipitation ofnanometric intermetallic phases like Ni4W andNiWmay be supposedas new and novel future trends.

• Tribological and mechanical properties. As discussed in previous sec-tions, Ni-W alloys have premium wear and mechanical features. Inthis regard, Hall-Petch breakdown was observed in nanocrystallinenickel-tungsten especially at grain sizes b10–15 nm. It seems thatthe Hall-Petch regime can be deep investigated still through alloyswith b10 nm grain size. Moreover, patterning the structure can benominated as a future direction. Layering the structure besides thesize and number of layers would illuminate new horizons in thisfield. In addition, tribological features and mechanical properties ofsuitable heat treated specimens may bring reliable and promising re-sults. Furthermore, developing new nanocomposite Ni-W based ma-terial with very high incorporated nanoparticles may be proposed asanother alternative of future trends.

• Thermal stability and high-temperature applications. Recently, the ther-mal stability of nanocrystalline Ni-W alloys has been investigatedthrough precise and nanoscopy approaches, introducing some raremetals element into these structures may enhance this ability. More-over, since nanotwinned metallic materials can exhibit unusual prop-erties, heavily nanotwinned Ni-W structures may be assigned forthermal stability investigations. Besides, developing diffusion barrierNi-W films for high-temperature coatings can be considered in thefirst and second stage turbine blades and vanes.

• Enhancing corrosion resistance. It may be possible to enhance the anti-corrosion properties of nickel-tungsten alloys by depositing soundand crack-free films, patterning the structure of incorporatinginhibition-effect/passive film forming nanoparticles. Investigation ofmultilayer structures with nanometric layer size is interesting regard-ing enhancing anti-corrosion properties of Ni-W alloys. Introducingsome alloying element for shifting the corrosion potential of thesecoatings, besides preserving the hardness can be regarded in the fu-ture trends. Besides, sinceMo can retard pitting corrosion of engineer-ing alloys (like 316 stainless steel), localized corrosion behavior of Ni-W-Mo coatings can be suggested as an interesting subject in this field.

• Hydrophobic features. Considering the main two factors influencingthe hydrophobicity of coatings including geometrical surface mor-phology and surface energy, it is possible to enlarge the contactangle by reducing the surface energy using some chemicals like

fluorocarbons, siloxane, and stearic acid and/or controlling of surfacemorphology and structure. Developing new electrolyte (such as alco-holic based ones) and using chemical agents likemyristic acid to treatthe Ni-W surface might be proposed as future trends in this field. Be-sides, taking into account the nucleation and growth principles, thetwo-step electrodeposition of nickel-tungsten can be applied to createmixedmicro and nano hierarchical morphologies, probably leading tohydrophobic surfaces.

• Barrier properties: Nanocrystalline Ni-W films exhibit effective barriercharacteristics either at low and moderately high temperature. Thispropertymay introduce these alloys as a reliable barrier against corro-sion, formation of intermetallics, migration of alloying elements, etc.in different industries. These alloys may be developed either as a dif-fusion barrier in the high-temperature application or effective barrierin electronic devices.

The implicit themes are but they are a few from among an immenselist of potentially motivated and novel directions for future research.With the recent prosperity of new findings, new characterization tech-niques and new researchers in this field, we look forward a rapid prog-ress on high quality amorphous and nanocrystalline nickel-tungstenbased alloys with premium and unique high temperature and mechan-ical properties.

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