Preparation and Properties of Textured Ni–W Coatings Electrodeposited on the Steel Surface from a Pyrophosphate Bath

Abstract

Ni–W alloys with a (2 2 0) or (1 1 1) preferred orientation growth and amorphous structure were prepared from a pyrophosphate bath using the electrodeposition method. Structure transformation can be the result of the bath temperature (T_b) and the concentration of sodium tungstate (C_W) in the bath. Increasing the T_b and C_W can change the crystal growth from (2 2 0) to (1 1 1). At a higher T_b and C_W, an amorphous Ni–W alloy can be obtained. The tungsten content in the coatings should be responsible for the structure change. The three textured Ni–W alloys with a (2 2 0) texture, (1 1 1) texture and amorphous structure were annealed at different temperatures ranging from 200 to 700 °C. The microhardness, corrosion resistance and HER of the as-deposited and annealed Ni–W alloys were comparatively studied. The results show that the microhardness of the amorphous Ni–W alloy is the highest and reaches 1028 HV after annealing at 400 °C. The (2 2 0)-textured Ni–W alloy has the best corrosion resistance, which is further improved after annealing, while the HER activity of the (1 1 1) textured Ni-W alloy is superior.

1. Introduction

A material’s texture has remarkable effects on its HER [1,2], corrosion [3,4,5], and other properties [5]. The properties of the HER in an alkaline solution are even changed by macroscopic tensile strains in nickel foils [2]. Xu et al. [1] reported that Ni with a (1 1 1) or (2 0 0) plane-preferred orientation was electrodeposited by changing the additives of NH₄Cl and (NH₄)₂SO₄. They found that the (111) preferential orientation promoted the electro-oxidation of NaBH₄, and the (200) and (220) planes, as preferred orientations, achieved optimal HER activity. After electrodeposition for 16 h, even the preferential orientation of electrodeposited Ni is changed, from the (111) and (200) planes to the (220) plane [4]. In the simple Watts bath, Nasirpouri [5] prepared Ni films with (1 1 1) and (1 0 0) preferential orientations by pulsed (PC) and pulsed reverse (PRC) current techniques, respectively. Eckold et al. [3] reported that the texture of electrodeposited Sn can be changed by increasing the current density, and a preferred orientation along the lattice planes (321) and (220) greatly enhances the corrosion resistance of the tin layer. In general, metals and alloys with different textures can be prepared by changing the additives and applied current density.

Compared with Ni coatings, Ni–W alloy coatings possess better wear resistance, corrosion resistance [6,7], and hydrogen evolution reaction (HER) activity [8,9,10,11]. They are also considered to be candidates to replace the environmentally hazardous hexavalent hard chromium coating [12]. Differently textured Ni–W alloys have been obtained from an ammonium-citrate bath by DC and PC electrodeposition [6,7]. However, in the bath, it is difficult to solve the problems of citrate decomposition, which leads to plating solution failure, decreased coating performance, and environmental pollution. As a non-cyanide complexant, pyrophosphate is not easily decomposed and consequently presents a longer lifespan. Ni–W crack-free coatings have been prepared from pyrophosphate baths [13,14].

Electrodeposited Ni–W alloys can be improved by heat treatment [15,16,17,18] and the introduction of other elements or compounds [17,18,19,20,21,22,23]. The synergistic effects and reinforcing phases of other elements and compounds can greatly improve the properties of Ni–W alloys. Heat treatment can cause changes in the metal microstructure, such as phase transformation, grain size, residual stress and cracks, consequently optimizing the properties of Ni–W alloy coatings by designing heat treatment processes [6].

In this work, a pyrophosphate bath was selected to prepare Ni–W alloys, and changing the concentration of sodium tungstate (C_W) and the bath temperature (T_b) could tune the structure of the Ni–W alloys, consequently forming a preferred orientation of the crystal face, e.g., (2 2 0), (1 1 1), or an amorphous structure. Heat treatment of these Ni–W alloys was carried out at 200 to 700 °C for 2 h in air. The microhardness, corrosion resistance, and HER of the as-deposited and annealed Ni–W alloys were comparatively studied.

2. Experimental Procedure

2.1. Preparation of Ni-W Alloy Coatings

The Ni–W coatings were electrodeposited from a pyrophosphate bath onto Q235 carbon steel disks with a diameter of 10 mm and a thickness of 2 mm. The chemical components of the Q235 steel are listed in

Table 1. The chemical components of the Q235 steel.

Component	C/%	Si/%	Mn/%	P/%	S/%	Fe/%
Q235 steel	0.12–0.2	<0.3	0.3–0.7	<0.045	<0.045	Remaining

.

Compared with the conventional ammonium-citrate system, the merit of the pyrophosphate bath is that there are no electro-productions resulting from the decomposition of organic chelating agents. To obtain different textured Ni–W alloys, a high bath temperature of 60 °C was also applied in this work. The hydrolysis of pyrophosphate is likely to occur at over 60 °C, so higher temperatures were not selected. The additive is only 2-butyne-1, 4-diol at 0.2 g L⁻¹. The bath composition and the electroplating parameters are listed in

Table 2. Bath compositions and plating conditions for textured Ni–W alloys.

Chemicals/Parameters	Values (g L−1)
NiSO4·6H2O	12
NaWO4·2H2O	30, 60
Na4P2O7·10H2O	40
H3PO4 (85%)	10
NH3·H2O (25%–28%)	~20
Bath temperature	40 °C, 60 °C
Current density	4–10 A dm−2
pH	8.5–9.0

.

Before electrodeposition, the carbon steel disks were ground using SiC waterproof sandpaper (CW-800), electrochemically degreased in 40 g L⁻¹ NaOH solution at a current of 10 A dm⁻² for 10 min, and then activated in 10% HCl solution for 10 s. Ni–W alloys were electrodeposited onto the pretreated carbon steel disks in a 500 mL pyrophosphate bath using a DC power supply (APS3005DM, ATTEN/Antaixin, Hangzhou Hamat Electronics Co., Ltd., Zhejiang, China). The bath temperature (T_b) was maintained at 40 or 60 °C. The electrodeposition time was 150 min to meet the requirements of heat treatment and microhardness testing. After electrodeposition, part of the carbon steel disks coated with Ni–W alloys were annealed at 200, 300, 400, 500, 600 and 700 °C for 2 h in a muffle furnace and then placed in a self-sealing bag for storage.

2.2. Characterization of Ni–W Coatings

The chemical composition of the Ni–W coatings was analyzed using an energy-dispersive X-ray spectrometer (EDS, EDX-8000, Shimadzu, Shanghai, China) in an air atmosphere with a collimator of 3 mm. The crystal structure was characterized by X-ray diffraction (XRD, D8 Advance, Bruker, Shanghai, China) in the 2θ range of 20 to 90° (Cu Kα radiation, λ = 0.15406 nm), and the grain size of the strongest diffraction peak was calculated by means of the Debye–Scherrer formula (D = 0.89 λ/ (β cosθ). The microscopic morphology was observed by field emission scanning electron microscope (FESEM, Nova450, FEI). The microhardness of the coatings was tested using a microhardness tester (HV-1000A, Laizhou Huayin Testing Instrument Co., Ltd., Laizhou, China) under an indentation load of 100 gf and a dwelling time of 10 at five different locations per specimen, and the average value was quoted as the hardness of the coating. The samples annealed above 400 °C were polished with W220, W7 and W3.5 metallographic sandpapers in turn because of surface oxidation. Thermo-analysis tests were carried out using a thermal analyzer (STA449F3, Netzsch, Shanghai, China) in air from room temperature to 800 °C at 10 °C min⁻¹. In order to carry out the thermo-analysis tests, the Ni–W alloys were electrodeposited onto stainless-steel disk and were then peeled from the substrate. The weight of all deposits for the thermo-analysis tests was about 10 mg.

The Q235 carbon steel disks coated with Ni–W alloy coatings were sealed with epoxy resin after welding with copper wire and polished with SiC waterproof sandpapers (CW-600, 800, 1000) to expose a sample surface of 0.785 cm². The electrochemical performance of the sample was tested using an electrochemical workstation (CHI 660E, Shanghai Chenhua, Shanghai, China). The saturated calomel electrode (SCE), platinum plate, and Fe coated with a Ni–W alloy electrode were directly used as a reference electrode (RE), counter electrode (CE) and working electrode (WE), respectively. The corrosion resistance of the Ni–W alloys was evaluated by polarization curves (Tafel) and electrochemical impedance spectroscopy (EIS) in a 3.5 wt. % NaCl solution. The Tafel was tested at a scan rate of 1 mV s⁻¹ with a relative open circuit potential (OCP) of ±250 mV, and the EIS was used with an AC signal amplitude of 5 mV at a frequency range of 0.01 Hz to 100 kHz. The behavior of HER was evaluated by linear sweep voltammetry (LSV) at a scan rate of 5 mV s⁻¹ in a 6 wt. % NaOH solution.

3. Results and Discussion

3.1. Structure of Three-Textured Ni–W Alloy Coatings

3.1.1. As-Deposited Ni–W Alloy Coatings

Figure 1 shows the XRD patterns of the Ni–W alloy coatings prepared at different current densities (4, 6, 8 and 10 A dm⁻²) and different concentrations of sodium tungstate (C_W) (30 or 60 g L⁻¹) at a T_b of 40 or 60 °C. It can be seen that the Ni–W coatings present diffraction peaks at 2θ of about 44°, 51°or 76°, corresponding to the three standard diffraction peaks of pure FCC nickel (JCPDS No: 04-0850). But, the peak position shifts towards a low angle, suggesting FCC lattice shrinkage because of the introduction of large atoms into the Ni matrix. Meanwhile, the higher diffraction angle compared to that of Ni₁₇W₃ (JCPDS No: 65-4828) implied a lower W content in the electrodeposited Ni–W alloys.

At a low C_W of 30 g L⁻¹ and a low T_b of 40 °C, a (2 2 0) textured Ni–W alloy was obtained, presenting an XRD peak at 2θ = 76°, as shown in Figure 1a. However, the Ni–W alloys obtained when C_W = 20 g L⁻¹ and T_b = 40 °C in Ref. [14] present the preferred (1 1 1) orientation. Compared with both baths, the difference is the applied additive, which is 2-butyne-1, 4-diol in this work and is a commercial additive for Ni–W alloys. This result indicates that the additive has an important effect on the growth of Ni–W alloy crystals.

When increasing the C_W from 30 to 60 g L⁻¹, the texture of a Ni–W alloy converts to the (1 1 1) orientation, as shown in Figure 1b. The intensity of the (1 1 1) peak becomes strong, and there is an increase in current density, implying the better crystallinity of the Ni–W alloy. However, although the corresponding Ni–W alloys reported in Ref. [14] exhibit the preferred (1 1 1) orientation, the peak intensity decreases with the increase in current density. The difference could result from the different additive. In the present system, the textured transformation from (2 2 0) to (1 1 1) should be attributed to the improved W content in the Ni–W alloys.

When increasing the T_b from 40 to 60 °C, the texture transformation from (2 2 0) to (1 1 1) can also achieve success, as shown in Figure 1c. When increasing the current density to 10 A dm⁻², a slight peak of (2 0 0) and a strong peak of (2 2 0) appear. It can be seen that the (1 1 1) peak slightly shifts towards a high angle, along with an increase in current density, suggesting lower W content in the Ni–W alloys.

When simultaneously increasing the C_W and T_b, the (1 1 1) peak becomes wide and low, as shown in Figure 1d, implying the formation of an amorphous structure. The higher C_W and T_b is beneficial to the W deposition during the electrodeposition of the Ni–W alloy. Under this condition, the (1 1 1) peak hardly rises with the current density; that is, current density has no effect on the crystallinity of Ni–W alloys.

In general, it can be considered that the structure change is mainly attributed to the increase in W in the Ni–W alloys or the selected additive.

The chemical composition of the Ni–W alloy coatings was measured by an energy-dispersive X-ray spectrometer in an air atmosphere with a collimator of 3 mm. The W and Ni content was calculated according to the characteristic peaks at 7.47 KeV for Ni and 9.67 KeV for W. The results are shown in Figure 2. At a low C_W of 30 g L⁻¹, as shown in Figure 2a, there is a decreasing trend for the W content along with an increase in the current density, while the trend is weak and even reversed at the high C_W of 60 g L⁻¹ (see Figure 2b). At the low C_W of 30 g L⁻¹, the W content increases along with the increase in T_b under the same current density conditions, as shown in Figure 2a. However, when increasing C_W to 60 g L⁻¹, it is difficult to increase the W content in Ni–W alloys by increasing T_b, especially over 50 °C. These results are consistent with the above XRD data. The diffusion control of the reduction in WO₄²⁻ ions and the limited deposition of tungsten induced by nickel is a reasonable explanation. Increasing the current density, the surface C_W becomes low because of the slow diffusion rate of species containing WO₄²⁻ ions, consequently resulting in the low W content in Ni–W alloys. When increasing the T_b, the surface C_W becomes high due to the increased diffusion rate from the T_b; therefore, the W content in the Ni–W alloys increases. At a high C_W of 60 g L⁻¹, the surface C_W is high and the deposition of tungsten is limited by nickel, so the W content in the Ni–W alloys is no longer changed by the T_b.

By controlling the C_W and the T_b, three textured Ni–W alloys could be obtained and were named W220, W111 and WAS with a (2 2 0) texture, (1 1 1) texture and amorphous structure, respectively. The electrodeposition parameters and the W content in the Ni–W alloys are listed in

Table 3. Electrodeposition parameters of three textured Ni–W alloys.

Name	W wt. %	Orientation	CW	Tb/°C	Current/A dm−2
W220	13.85	(2 2 0)	30	40	8
W111	29.80	(1 1 1)	60	40
WAS	36.75	Amorphous	60	60

.

Figure 3 shows the SEM images of the three textured Ni–W alloy coatings. It can be seen from Figure 3a,b that crystal particles with sizes between 20 and 200 nm gathered together. However, many micron particles can be observed from the surface of the Ni–W alloy coating under the without-additives condition [14]. The result indicates that the crystal size of the Ni–W alloys can be markedly refined by the additives of 2-butyne-1, 4-diol. But, for the WAS sample seen in Figure 3c, the crystal grains cannot be observed from the image with the same magnification (50,000×), and the surface is compact and flat, being related to the amorphous structure. As shown in Figure 3d, the thickness is estimated to be 66, 68 and 81 μm for the W220, W111 and WAS samples, respectively, judging by the cross-sectional view.

3.1.2. Annealed Ni–W Alloy Coatings

Figure 4 shows the XRD pattern evolution of the W220, W111 and WAS Ni–W alloys, along with the increase in annealed temperature. From Figure 4a, it is found that below 700 °C, the (2 2 0) preferred orientation became stronger and stronger when increasing the annealed temperature, while other diffraction peaks became slightly stronger. During heat treatment, the (2 2 0) crystal plane grew rapidly, and the (2 2 0) textured structure could be retained until a temperature of 600 °C was reached. At 700 °C, the Ni–W alloy was oxidized, and new diffraction peaks appear in the magnified curve [24].

As shown in Figure 4b, there is the (111) preferred texture in the W111 sample until a temperature of 500 °C is reached. During heat treatment below 500 °C, there is almost no growth of the (111) crystal surface, while the (2 0 0) and (2 2 0) plane distinguishably grew. The oxidization of Ni–W also happened at 700 °C.

For the WAS sample, the broad and weak (1 1 1) peak becomes strong and narrow with an increasing annealed temperature, suggesting that fine grains were formed during the transformation from the amorphous to the nanocrystalline structure. The (2 0 0) and (2 2 0) peaks are weaker than those of the W111 sample, implying that their growth is suppressed. At 600 °C, the strongest (1 1 1) peak suggests that the transformation from amorphous to crystalline was completed.

The differential scanning calorimetry (DSC) and thermogravimetry (TG) curves of the W220, W111 and WAS samples are shown in Figure 5. It can be seen from the DSC curves that all samples present an exothermic process, which accelerates at about 500 °C. The exothermic process is related to crystal growth, so the preferential diffraction peaks become strong and are strongest after 500 °C, as shown in Figure 4. The severe oxidation at 700 °C made the surface of the Ni–W alloys black (see the inset in Figure 5) and undermined the diffraction peaks of the Ni–W alloys. As a result, the strongest diffraction peaks appear at 600 °C, as shown in Figure 4.

The TG curves present the different shapes. It can be seen from Figure 5 that for the W220 sample, there is almost no change in weight below 335 °C, and then, the weight increases. At the same time, the surface of W220 appears red and even brown at 400 °C, as shown in the insets of Figure 5a.

For the W111 and WAS samples, the weigh firstly decreases, plateaus, and then increases along with the temperature, as shown in Figure 5b,c. The weight loss may result from airborne desorption because the (111) crystal face possesses higher activation than the (220) crystal face. For the WAS sample, more weight loss is observed, and the formation of volatile tungsten oxide on the surface is a reasonable explanation. The initial stage of the plateau corresponds to the start of oxidation because of the appearance of surface color at about 400 °C. At 400 °C, the color becomes gradually light for the W220, W111 and WAS samples, which is consistent with the increase in W in the alloys. At over 600 °C, the rapid increase in weight and the deepening in color indicate a faster oxidation rate.

Based on the (2 2 0)’s full width at half maxima in Figure 4a and the (1 1 1) peak in Figure 4b,c, the grain size of the alloys were calculated using the Debye–Scherrer formula. The variation in grain size of the W220, W111, and WAS Ni–W alloys with temperature is displayed in Figure 6. The results show that the grain size of the as-deposited W220 is the largest (10.78 nm), and that of WAS is the smallest (9.08 nm). It can be concluded that the grain size of the Ni–W alloy decreases with the increase in W content. And, the grain size of the Ni–W alloy 2 h after annealing at 200 and 300 °C hardly changes. As the heat treatment temperature increases, the grains of the coatings grew to 20.60~39.82 nm at 600 °C. At 700 °C, the grain of WAS further coarsened to 30.82 nm, while W220 and W111 became small because of the formation of the new oxidation phase, presenting diffraction peaks at 2θ of 23.9 to 36.3°.

3.2. Microhardness

The microhardness of the three textured Ni–W alloy coatings is shown in Figure 7. The microhardness of the as-deposited W220, W111 and WAS samples ranged from 550 to 600 HV, being similar to the reported value of 460–740 HV [17]. It can be seen that the microhardness of WAS with the highest W content is highest within the entire annealed temperature range because of the solid solution strengthening effect of W [25]. Below 400 °C, the microhardness of the three textured Ni–W alloy coatings rapidly improved after heat treatment, being attributed to the elimination of residual stress and grain refinement crystallization [6,17].

At 400 °C, the microhardness of WAS reached 1028 HV. However, the microhardness for W220 was only 760 HV because its W content is lowest (see Table 2). Over 400 °C, the microhardness decreases along with the annealed temperature because of the coarsening crystal grains.

3.3. Corrosion Resistance

3.3.1. Tafel Curves

The corrosion resistance of the as-deposited and annealed Ni–W alloy coatings was characterized by Tafel in a 3.5 wt.% NaCl solution. For comparison, the matrix Q235 carbon steel disk was also used.

Figure 8 shows the Tafel curves of the three textured Ni–W alloy coatings of W220, W111 and WAS that were as-deposited and annealed at different temperatures in air for 2 h. And, the corrosion potential (E_corr) and corrosion current density (I_corr) were obtained by the Tafel extrapolation method and are given in

Table 4. Corrosion resistance of three textured Ni–W alloy coatings.

Sample	T/°C	Icorr/uAcm−2	Ecorr/mV vs. SHE	ΔEcorr/mV	IE/%
Fe	-	5.71	−370	0	-
W220	As-deposited	13.2	−205	165	−131.2
	200	5.77	−174	196	−1.1
	400	2.55	−64	306	55.3
	600	0.81	−41	329	85.8
W111	As-deposited	16.27	−231	139	−184.9
	200	10.93	−207	163	−91.4
	400	2.00	−127	243	65.0
	600	3.22	−68	302	43.6
WAS	As-deposited	16.82	−258	112	−194.6
	200	11.49	−236	134	−101.2
	400	11.41	−223	147	−99.8
	600	2.56	−96	274	55.2

. It can be seen that the E_corr saw a positive shift after these three textured Ni–W alloy coatings were electrodeposited onto the Q235 carbon steel disks, implying improved corrosion resistance [24]. The more the annealed temperature elevated, the more the E_corr for all the samples is positively shifted. For the WAS sample, its E_corr is significantly positively shifted after heat treatment at 600 °C, likely due to the completion of the structural transformation from amorphous to crystalline and the elimination of internal stress, which results in cracks in Ni–W obtained in an ammonium-citrate bath [17]. The ΔE_corr was calculated according to the difference between the samples and Fe blank electrode. The ΔE_corr is 165 mV, 139 mV and 112 mV for the as-deposited W220, W111 and WAS samples, respectively. That is, the value of ΔE_corr gradually decreases along with the increase in W content in Ni–W alloys. It was found that I_corr becomes high after coating Ni–W alloys, possibly because the galvanic coupling of Fe and Ni–W accelerate the corrosion of Fe in porous coatings. For every sample, the I_corr decreases along with the increase in annealed temperature. In general, the I_corr of the W220 sample is the lowest, which suggests that W220 possesses the best corrosion resistance. Alimadadi also reported that the corrosion resistance of single-phase Ni–W is superior to amorphous and dual-phase coated layers because of cracks or internal stress [26].

Improvement of corrosion resistance can be analyzed by the inhibition efficiency (IE, %), which was calculated according to the following equation (the results are shown in Table 4).

$$\mathrm{I}\mathrm{E}=\frac{{\mathrm{I}}_{corr, Fe}-{\mathrm{I}}_{corr, NiW}}{{\mathrm{I}}_{corr, Fe}}\times 100\mathrm{\%}$$

It can be seen that the IE decreased after the electrodeposition of the Ni–W alloy onto Q235 carbon steels, but IE can be increased by means of the heart treatment, and the biggest IEs are 85.8% for the W220 sample, 65.0% for the W111 sample and 55.2% for the WAS sample. As shown in Figure 5, the oxidation behavior occurs at temperatures of 400 °C or higher, which maybe decrease the porosity of the coating, and consequently, the results indicate that the W220 sample presents the best corrosion resistance.

3.3.2. EIS

The corrosion resistance of the Ni–W alloy coatings was also estimated by EIS spectra measured in the 3.5 wt. % NaCl solution, and the results are shown in Figure 9. All the Nyquist curves present a flat semicircle, suggesting that the Fe, as-deposited, and annealed Ni–W alloy coatings had undergone the same corrosion reaction process. However, the semicircular size of the as-deposited Ni–W alloy coatings is much larger than that of Fe, indicating the corrosion resistance is improved after electrodepositing Ni–W alloy coatings onto the Fe substrate. The semicircular arc became extraordinarily larger after heat treatment. Therefore, corrosion resistance can be significantly improved by annealing after electroplating.

Fitting through an equivalent circuit diagram shown in Figure 9d, the results are shown in

Table 5. EIS Fitting results of three textured Ni–W alloys according to the equivalent circuit shown in Figure 9d.

Sample	T/°C	Rs/Ω cm2	CPE-T/uF cm−2	CPE-P	Rct/Ω cm2	IERct/%
Fe	-	5.98	1264	0.753	1059	-
W220	As-deposited	6.56	419	0.69	10,743	90.1
	200	5.26	223	0.74	12,065	91.2
	400	6.12	110	0.78	18,012	94.1
	600	5.85	78	0.89	43,267	97.5
W111	As-deposited	6.284	106	0.72	5826	81.8
	200	6.748	137	0.77	6511	83.7
	400	5.07	49	0.88	6325	83.2
	600	6.043	46	0.88	28,178	96.2
WAS	As-deposited	5.63	530	0.78	2918	63.7
	200	4.525	219	0.77	4751	77.7
	400	5.821	137	0.78	8083	86.9
	600	5.91	4	0.663	30,154	96.5

. Here, R_S is solution resistance, R_ct is charge transfer resistance, and CPE₁ is the double layer capacitance involved in the equivalent circuit. It can be seen that the R_ct value increases from 1059 Ω cm² for the Fe electrode to 10,743 Ω cm² for the W220 sample, 5826 Ω cm² for the W111 sample, and 2918 Ω cm² for the WAS sample. The R_ct value of the coatings was dramatically increased, along with annealed temperature. The maximum value of R_ct is 43,267 Ω cm² for the W220 sample, 28,178 Ω cm² for the W111 sample, and 30,154 Ω cm² for the WAS sample. The inhibition efficiencies were also calculated by R_ct according to the following equation for IE_Rct, and the results are shown in Table 5.

$${\mathrm{I}\mathrm{E}}_{Rct}=\frac{{\mathrm{R}}_{ct, Ni-W}-{\mathrm{R}}_{ct, Fe}}{R_{ct, Ni-W}}\times 100\mathrm{\%}$$

As expected, for all samples, the IE_Rct values are all positive and increase with annealed temperature. The IE_Rct of the as-deposited W220, W111 and WAS is 90.1%, 81.8% and 63.7%, respectively. Under the same condition, the WAS sample presents the biggest IE_Rct, indicating the best corrosion resistance, which is consistent with that of the Tafel curves. For the three samples, the biggest value of IE_Rct is 97.5% for W220, 96.2% for W111 and 96.5% for WAS.

3.4. HER

The electrocatalytic HER performance of the Ni–W alloy coatings was evaluated by LSV electrochemical technology in a 6 wt. % NaOH solution, and the results are shown in Figure 10. Compared with the Fe electrode, the W220, W111 and WAS alloys all showed improved electrocatalytic activity. With a driving current density of 100 mA cm⁻², the HER potential required was −1.400 V for W220, −1.387 V for W111 and −1.413 V for WAS. As we know, the lower the potential is, the better the HER activity is. This result indicates that the HER activity of the (1 1 1) textured Ni–W alloy is superior to the (2 2 0) textured and amorphous samples. The LSV curves of the three samples annealed at 200, 400 and 600 °C are between that of Fe and the as-deposited Ni–W alloys; that is, the high HER overpotential is presented after heart treatment. Therefore, it is not beneficial to promote the HER property by means of heart treatment after electrodeposition.

4. Conclusions

In summary, Ni–W alloys with a (2 2 0) or (1 1 1) preferred orientation growth and amorphous structure were prepared by electrodeposition from a pyrophosphate bath. Low C_W and T_b are conducive to the formation of a (220) plane preferential orientation, while on the contrary, a (111) orientation grew. The amorphous Ni–W alloys (WAS) with higher W content had the highest microhardness, being up to 1028 HV at 400 °C. The as-deposited and annealed Ni–W alloys with a (2 2 0) texture had better corrosion resistance, exhibiting a ΔE_corr of 329 mV, an I_corr of 0.81 uA cm⁻² and an R_ct of 43,267 Ω cm². The initial HER overpotential of the (1 1 1) textured Ni–W alloy was the smallest—560 mV at 10 mA cm⁻².