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Electrochemical evaluation of corrosion and tribocorrosion behaviour of amorphous and nanocrystalline cobalt–tungsten electrodeposited coatings
Materials Chemistry and Physics xxx (2014) 1e10
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Electrochemical evaluation of corrosion and tribocorrosion behaviour of amorphous and nanocrystalline cobaltetungsten electrodeposited coatings N. Fathollahzade, K. Raeissi* Department of Materials Engineering, Isfahan University of Technology, Isfahan 84156-83111, Iran
h i g h l i g h t s Mass-transport effect had higher proportion in tribocorrosion of CoeW coatings. The major electrochemical-wear degradation was for the nanocrystalline coating. The higher proportion of wear accelerated corrosion was for the amorphous coating. Superficial microcracks were formed near scars due to the coatings brittleness.
a r t i c l e i n f o
a b s t r a c t
Article history: Received 4 November 2013 Received in revised form 4 May 2014 Accepted 9 July 2014 Available online xxx
Amorphous and nanocrystalline CoeW coatings were electrodeposited on copper substrates from a citrateeammonia bath. The coatings showed nodular surface morphologies, but a microcrack network was detected in the amorphous coating. However, a better corrosion resistance was achieved for the amorphous coating. During sliding under open circuit potential (OCP) condition, the potential of amorphous coating gradually became more active probably due to the widening of wear scar, and thus expansion of active area. The amorphous coatings showed a higher volume loss at OCP probably due to its lower microhardness. In anodic sliding, a sharp increase in current density was observed due to mass transport and depassivation effects. In all sliding conditions, the proportion of mass transport was higher than wear accelerated corrosion, which implied that the dissolution reaction of the coatings was mainly a mass-transport controlled process. The results also showed that the effect of sliding on degradation is more intense for the nanocrystalline coating. For both coatings, the formation of the superficial microcracks in the vicinity of wear scars indicating on a surface fatigue wear mechanism. © 2014 Elsevier B.V. All rights reserved.
Keywords: Coatings Amorphous materials Nanostructures Corrosion Tribology
1. Introduction Cobalt and nickel based alloys containing tungsten have the potential to become promising replacements for hard chromium coatings deposited from toxic, carcinogenic and environmentally unfriendly hexavalent chromium baths [1,2]. Wang et al. [3] reported improved tribological properties of electrodeposited Co over Ni coatings with hexagonal close packed (hcp) and face centred cubic (fcc) structures, respectively, and attributed this result to the higher resistance of hcp structure to adhesion interactions during the wear process. Higher W content in bath, smaller grain size [4]
* Corresponding author. Fax: þ98 3113912752. E-mail address: [email protected] (K. Raeissi).
and strong hcp (100) crystal orientation [5] improve the hardness of the nanocrystalline CoeW coatings which is higher than that of the amorphous structures [6]. Su et al. [5] obtained a microhardness range of 800e1000 HV for nanocrystalline CoeW alloys electrodeposited under various conditions, while CoeW coatings produced by Ghaferi et al. [1] had a microhardness of 280e620 HV. Capel et al. [7] reported an increase in hardness (from 540 to 1200 HV) and a significant decrease in corrosion resistance by heat treatment. Some authors found a lower rate of wear for CoeW electrodeposit compared to pure cobalt electrodeposit [8] on account of the incorporations of W, and thus the increased hardness [9]. Tribocorrosion is the degradation of materials due to the simultaneous effects of electrochemical corrosion and mechanical wear in a tribological contact [10]. The total material removal rate
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Please cite this article in press as: N. Fathollahzade, K. Raeissi, Electrochemical evaluation of corrosion and tribocorrosion behaviour of amorphous and nanocrystalline cobaltetungsten electrodeposited coatings, Materials Chemistry and Physics (2014), http://dx.doi.org/ 10.1016/j.matchemphys.2014.07.013
N. Fathollahzade, K. Raeissi / Materials Chemistry and Physics xxx (2014) 1e10
differs from sum of the wear rate measured in the absence of corrosion and the corrosion rate observed in absence of wear [11]. For example, passive coatings would subject to active corrosion due to the annihilation of protective passive film in wet sliding contact, when the frictional movement in corrosive medium is continuous [12]. Regardless of the considerable corrosion data on cobalt and its alloys available in the literature [1,13e21], the tribocorrosion behaviour of them have been less investigated. Recently, Ghosh and Celis [22] reported that the open circuit potential (OCP) of CoeW alloy in 0.5 M NaCl electrolyte became nobler during the sliding and remained at that nobler value even after the end of sliding test. However, the OCP decreased progressively and reached the value recorded before starting the sliding test until the corundum ball was unloaded and removed from the sample surface [22]. These researchers [22] attributed this behaviour to the main role of tungsten in CoeW coatings which forms WO3 and catalyses the oxygen reduction reaction (ORR). Nevertheless, Liu et al. [23] suggested WO3 as a possible support material for monolayer Pt ORR electrocatalysts in acid electrolytes and have no idea about catalytic effect of WO3 on ORR. In this study, corrosion and tribocorrosion behaviours of amorphous and nanocrystalline CoeW coatings electrodeposited from a citrate-ammonia bath are investigated. The evolution of OCP values, anodic and cathodic current readings and SEM images taken from wear scars are applied to determine tribocorrosion behaviour of the coatings. The tribocorrosion resistance of the coatings were also estimated by calculating the volume loss in worn track using line profilometry.
Fig. 1. Schematic representation of reciprocating ball-on-plate tribometre. (1) DC motor, (2) specimen (working electrode), (3) counter electrode, (4) reference electrode, (5) load cells, (6) ball and its holder, (7) spring, (8) loading bolt (9) connecting rod, (10) potentiostat.
electromagnetic (eddy-current) method. The grain size of the coatings was calculated from X-ray diffraction (XRD) patterns obtained by an X-ray diffractometer (model X'pert Philips) with Cu Ka radiation (l ¼ 1.5418 Å), using the Scherrer equation. For the curve fitting analysis required for Full-Width at Half Maximum (FWHM) determination, the four-parameter Gaussian function was employed [24]. The instrumental line broadening was measured using a silicon peak corrected by following equation:
b/bexp ¼ 1 (bins/bexp)2
where b is the modified FWHM, bexp is the calculated FWHM of experiment specimen and bins is the FWHM of monocrystalline silicon. A Philips CM120 transmission electron microscope (TEM) with 0.037 Å wavelength and 100 kV voltage was used to detect the grain size. TEM sample preparation was performed by scratching the coating from substrate using a surgical blade and the resulted powders were dispersed in alcohol to choose a suitable particle for analysis.
(101)
Cu
(b)
Table 1 Composition of the bath for CoeW electrodeposition. Substance
Amount (g l1)
CoSO4$7H2O Na2WO4$2H2O NH4CL NaBr Na3C6H5O7$2H2O
40.24 51.22 26.74 14.34 178.58
Cu
(002)
The copper used as substrate was a disk of 0.9 cm2 surface area. The substrate surface was abraded with wet SiC papers up to grit size no. 1200 and then anodically polished for 15 min in a bath containing 80 vol.-% concentrated phosphoric acid and 20 vol.-% distilled water at 1.4 V, 25 C and 700 rpm. The cathode was a large plate of stainless steel. Immediately after washing with distilled water, the specimen (as the working electrode) was introduced into a cell containing electrodeposition solution with the composition shown in Table 1. The pH of the bath was adjusted at 7.5 using dilute NaOH and H2SO4 solutions. A platinum plate was applied as a counter electrode. The electrodeposition process was performed using a digital coulometer (BHP 2050). The amorphous and nanocrystalline CoeW alloys were electrodeposited at deposition current densities of 70 and 15 mA cm2 and temperatures of 45 and 25 C, respectively. The amorphous coatings were electrodeposited at 110 rpm, whereas nanocrystalline coatings were obtained at stationary conditions. The times for amorphous and nanocrystalline CoeW electrodeposition were 500 and 2330 s, respectively. Philips XL30 scanning electron microscope (SEM) was used to study the surface morphology and topography of the coatings before and after tribocorrosion experiments. Elemental compositions of the coatings were determined using an electron dispersive spectroscope (EDS) coupled with a Seron Technology AIS 2300C. Thickness of the coatings was measured about 15 mm using the
(1)
(100)
2. Materials and experimental methods
Intensity
2
(a) 30
35
40
45
50
55
60
2θ (degree) Fig. 2. XRD patterns of electrodeposited (a) amorphous (with 48 wt-% W) and (b) nanocrystalline (with 35 wt-% W) CoeW coatings.
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N. Fathollahzade, K. Raeissi / Materials Chemistry and Physics xxx (2014) 1e10
3
were recorded using an EG&G potentiostat/galvanostat (model 263A). Tribocorrosion tests under anodic and cathodic polarisation were done by applying constant potentials of þ0.1 V/OCP and 0.7 V/Ag/AgCl, respectively. The specimen was relaxed for 1200 and 1000 s before and after sliding, respectively, then the ball was unloaded and the recording were continued further for 1000 s. The speed and stock length of a complete cycle of reciprocating movement on the specimen surface were 1 cm s1 and 1 cm, respectively. After sliding, line profiles of wear scars were drawn using a surface profilometer (Mitutoyo Surftest SJ-301), in order to determine the shape and depth of the wear scars, and calculate the volume loss of specimen. In order for simulation of the fluid dynamics as in real tribocorrosion experiments, the ball was placed at a distance of about 0.5 mm above the surface of the working electrode and ball motion was done without any contact with the specimen surface under OCP and anodic/cathodic polarisation conditions. Each measurement was repeated for three times. 3. Results and discussion Fig. 3. Bright field of TEM image from nanocrystalline coating.
Microhardness of the coatings was determined using Buehler Micromet 5101 by applying a constant load of 25 g for 10 s. The calculated value of microhardness was the average number of five random points on the surface of each coating. Each measurement was repeated for three times. All electrochemical experiments were done using platinum as the auxiliary (counter) electrode, saturated Ag/AgCl as the reference electrode (with offset potential of 0.197 V/SHE) and the coated specimen as the working electrode. The testing electrolyte was 3.5 wt-% NaCl solution applied at room temperature in stationary condition. The tests were run after 60 min immersion of the working electrode. In all representations, cathodic currents have been defined as negative values. Five specimens were taken for each electrochemical experiment. Potentiodynamic polarisation measurements were carried out using a Princeton Applied Research (PARSTAT) potentiostat model 2273 at a sweep rate of 1 mV s1. Electrochemical impedance spectroscopy (EIS) measurements were run at OCP, before and after sliding, using an EG&G AC responser (model 1025) coupled with an EG&G potentiostat/galvanostat (model 263A). The voltage amplitude and the frequency range were 10 mV and 100 kHze10 mHz, respectively. Tribocorrosion experiments were performed using a reciprocating ball-on-plate tribometre described in Fig. 1. The working electrode surface was rubbed with a SiC ball under the load of 2 N for 1800 s at 1 Hz. The OCP values and measured current readings
3.1. Characterisation of CoeW coatings XRD patterns of the electrodeposited amorphous and nanocrystalline CoeW coatings are shown in Fig. 2. The pattern of nanocrystalline coating shows three characteristic peaks around 2q of 41.3, 44.3 and 47 which are very close to those in JCPDS 05-0727 attributed to (100), (002) and (101) planes of hcp Co, respectively. However, the shifts in peak positions of hcp Co to lower 2q is observed in Fig. 2 and related to the formation of a supersaturated solid solution of W in Co. In this case, since Co atoms in the lattice are substituted by bigger W atoms, an increase in the average closest distance between two adjacent atoms should be expected. On the other hand, a broad peak around 2q range of 40e50 is observed in the pattern of amorphous coating that confirms the amorphous structure of the alloy. Since, the characteristic peaks of W and Co3W are absent in Fig. 2, thus complete dissolution of tungsten in cobalt has occurred. For both coatings, the characteristic peaks of copper substrate are evidenced around 2q of 43.3 and 50.4 belonging to (111) and (200) planes of fcc copper which are consistent with cubic Cu lines in PDF 04-0836. Using the Full-Width at Half Maximum (FWHM) of Co (100) peak and the Sherrer equation, the grain size of nanocrystalline CoeW coating was estimated to be around 15 nm. The TEM result also confirms a same average grain size as is seen in Fig. 3. A nodular surface morphology can be observed in SEM images of the amorphous and nanocrystalline CoeW electrodeposited coatings (Fig. 4). Furthermore, a microcrack network is evidenced in the amorphous CoeW coating. The formation of the microcrack network could be related to the extensive hydrogen evolution
Fig. 4. SEM images of the area used for EDS analysis of (a) amorphous and (b) nanocrystalline CoeW coating. The higher magnification pics are also inserted.
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N. Fathollahzade, K. Raeissi / Materials Chemistry and Physics xxx (2014) 1e10 A - BNC
A - NS
A - SC
N - SC
N - NS
S - NS
stationary and ball motion conditions are shown in Fig. 5. Data extracted from Fig. 5 using Tafel extrapolation method are listed in Table 2. The polarisation resistance, R(p), is calculated using the SterneGeary equation [27]:
N - BNC
0.01
Current density (A cm-2)
0.001
R(p) ¼ (babc)/(2.3ic(ba þ bc))
0.0001
0.00001
0.000001
0.0000001 -0.9
-0.8
-0.7
-0.6
-0.5
-0.4
-0.3
-0.2
-0.1
0
Potential (V/Ag/AgCl) Fig. 5. Potentiodynamic polarisation curves (E vs. log i) of Cu substrate (S) and the amorphous (A) and nanocrystalline (N) CoeW coatings under different sliding condition in 3.5 wt-% NaCl solution. (NS: no sliding, SC: sliding with contact, BNC: ball motion with no contact).
coming up by increasing the current density and temperature of deposition process [1]. However, some authors have attributed this to the internal stresses built up by the presence of a high percentage of tungsten in the amorphous coating [25]. EDS analysis revealed that the amorphous and nanocrystalline CoeW electrodeposits had about 48 and 35 wt-% tungsten, respectively. It is reported that the transition from nanocrystalline phase to amorphous depends on tungsten content and can occur when it exceeds a critical value (about 45 wt-%) making the disordered state energetically more favourable [26]. Despite of the higher content of tungsten in amorphous coating, nanocrystalline coating had higher microhardness value. Vickers microhardness of the amorphous and nanocrystalline CoeW coatings were about 290 and 360 HV, respectively, which are lower than those reported by some authors [5,7,9]. It has been reported that the hardness of CoeW alloy coatings greatly depends on the content of W, the grain size, crystal orientation and phase structure [5,9] which are influenced by electrodeposition conditions such as bath composition and deposition current density. Incorporations of W in Co up to ~20 wt-% [5] causes a significant increase in the hardness of CoeW alloys owing to the decrease of the grain size, however, larger incorporation of W leads to a less strong hcp (100) orientation [5] which counteracts with the decrease of the grain size and thus decreases the hardness. 3.2. Potentiodynamic polarisation measurements Potentiodynamic polarisation plots (E vs. log i) for Cu substrate, the amorphous and nanocrystalline CoeW electrodeposits at
(2)
where ba and bc are Tafel constants for the anodic and cathodic reactions, respectively, and ic is the corrosion current density. As is seen in Table 2, the coatings which are under ball motion (with or without contact with the specimen surface) show lower R(p) with E(corr) values shifted to the more active values. This is more intense in contact (sliding) condition. However, the changes observed in OCP values were negligible (Table 2). According to R(p) values, the amorphous CoeW coating showed better corrosion resistance than the nanocrystalline one. In each case, E(corr) of the amorphous coating was nobler than that of the nanocrystalline one, whereas the nanocrystalline coating showed nobler OCP (Table 2). The main reason behind the difference exists between OCP and E(corr) of the coatings (Table 2) is more likely the reduction of passive layer whilst drawing the cathodic branch in potentiodynamic polarisation curves, which causes a shift in mixed potential (E(corr)) to a more negative (active) value. In the absence (or weakness) of passive layer on the coatings surface, the amorphous coating should act nobler than the nanocrystalline one at least due to the presence of microcracks, and hence interference of the nobler Cu substrate in measurements. In the presence of passive layer, which may plug most of the microcracks in amorphous coating (Fig. 6a), the influence of Cu substrate is diminished in OCP measurements. Fig. 6 shows SEM images of the amorphous coating in secondary electrons (SE) and back scattered electrons (BSE) modes after 60 min immersion in NaCl solution, just before starting the potentiodynamic polarisation readings. The micrographs confirm that the microcracks are plugged by corrosion products which are seen darker than the matrix in BSE mode (Fig. 6b), because of their lighter weight. Although, the tungsten content and surface morphology of a coating affect the corrosion resistance, but, the amorphisation of coating which happens at higher tungsten content is more likely responsible for the higher corrosion resistance. Generally, the amorphous materials are more corrosion resistant than the crystalline ones on account of the absence of crystallographical defects such as grain boundaries. In spite of the presence of microcracks in amorphous CoeW coating (Fig. 4a), it has still shown a lower corrosion rate compared to the nanocrystalline coating. As was discussed, this is probably because of the microcracks plugging by corrosion products (Fig. 6). Certainly, the microcracks plugging could prevent the build-up of galvanic corrosion cells, and thus eliminate the negative effect of the microcracks on corrosion resistance of the amorphous coating (prevents crevice corrosion).
Table 2 Data extracted from potentiodynamic polarisation plots in Fig. 5. (A: amorphous, N: nanocrystalline, S: substrate). Parameter
Value No motions
E(corr) (mV) i(corr) (mA cm2) ba (V dec1) jbcj (V dec1) OCP (mV) R(p) (U cm2)
Sliding
Ball motion with no touching the specimen
(A)
(N)
(S)
(A)
(N)
(A)
(N)
680 6 0.13 0.06 530 2975
690 8 0.17 0.05 520 2100
170 12 0.07 0.27 150 2015
710 14 0.08 0.08 533 1240
720 18 0.09 0.09 522 1090
700 12 0.09 0.07 532 1425
710 15 0.09 0.09 521 1305
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Fig. 6. SEM images of amorphous CoeW electrodeposited coating after 60 min immersion in 3.5 wt-% NaCl solution (before potentiodynamic polarisation test) obtained by detection of (a) secondary (SE) and (b) back-scatter (BSE) electrons.
3.3. Tribocorrosion experiments at OCP
(a)
-520
Potential (mV/Ag/AgCl)
OCP values achieved during the tribocorrosion experiments are related not only to the electrochemical condition of working electrode [22], but also depend on ratio of worn to unworn areas and mass transport condition built up during the sliding [28]. In passive metals, cyclic repassivation/depassivation is also an agent which determines OCP in tribocorrosion experiments. Fig. 7 shows OCP values of the amorphous and nanocrystalline coatings before, during and after ball motion (with and without contact with the working electrode surface). As is seen in Fig. 7, the ball motion (in both contact conditions) causes a significant negative drop in OCP for both coatings. The ball motion alone would enhance the dissolution of the passive film covering the coatings surface [28], if its dissolution were under mass-transport control [29]. It is well known that the thickness of passive films is determined by the
-530
Unloading the ball
-540
(2)
Start sliding
-550
(1)
-560
End sliding
-570 -580 0
Potential (mV/Ag/AgCl)
(b)
1000
2000 3000 Time (s)
4000
5000
-510 Unloading the ball
-520
(2)
-530 Start sliding
-540 -550
End sliding
-560
(1)
-570 0
1000
2000 3000 Time (s)
4000
5000
Fig. 7. Evolution of OCP values of (a) amorphous and (b) nanocrystalline CoeW coatings before, after and during tribocorrosion experiments (1) with and (2) without touching the specimen surface in 3.5 wt-% NaCl solution.
competitive effects of applied potential and film dissolution rate [30]. Thus, by increasing the dissolution rate of passive film due to the mass transport effect (as will be evidenced in Section 3.6), the film becomes thinner and OCP drops to negative values as is seen in Fig. 7. In the case of sliding, OCP is determined not only by the mass transport, but also by some other effects such as: (1) degree of annihilation of passive layer (depassivation) by sliding and extent of active area in worn track, (2) formation of microcracks in the vicinity of wear scar due to the brittleness of coatings, which may enhance corrosion. During the sliding, OCP value gradually became more negative for the amorphous coating (Fig. 7a), whilst this value slightly increased just by ball motion (with no contact). For the nanocrystalline coating (Fig. 7b), OCP slightly increased during the ball motion in both contact conditions. At the end of sliding, OCP values became nobler but they did not approach their initial values. Finally, a significant decrease in OCPs could be observed after unloading the ball from the specimen surface. This decrease is more likely due to the exposure of a very small surface beneath the ball to the solution, which acts more active than other areas and exaggerates the galvanic corrosion. According to Fig. 7a, OCP value of the amorphous coating gradually becomes more active when sliding is going on. This is more probably due to the higher trend of increase in wear scar width of the amorphous coating during the sliding, which surly extends the active area. This more likely arises due to the lower microhardness of the amorphous coating compared with the nanocrystalline one. For the nanocrystalline coating, the variation of OCP during the sliding is very similar to that observed just by ball motion (Fig. 7b) denoting that the total active area in wear scar is likely remained near constant during the sliding. This more possibly arises owing to the higher microhardness of the nanocrystalline coatings. At the end of sliding, restoration of the passive film took place on the worn areas, and hence OCP values gradually became nobler. 3.4. Tribocorrosion experiments under anodic and cathodic polarisation Fig. 8 shows current density readings under anodic polarisation. At the start of applying the anodic potential, the current density of both amorphous and nanocrystalline coatings decreased markedly owing to the thickening of passive layer. It is well known that the electrochemical conditions required for forming passive films are more pronounced at the anodic polarisation than OCP condition. Moreover, under anodic polarisation, the electron extraction force is so great and hence electrons are extracted not only from the conduction band but also from the valence band [31]. In this case, an inversion layer is formed in the space-charge region and the
Please cite this article in press as: N. Fathollahzade, K. Raeissi, Electrochemical evaluation of corrosion and tribocorrosion behaviour of amorphous and nanocrystalline cobaltetungsten electrodeposited coatings, Materials Chemistry and Physics (2014), http://dx.doi.org/ 10.1016/j.matchemphys.2014.07.013
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N. Fathollahzade, K. Raeissi / Materials Chemistry and Physics xxx (2014) 1e10
(a)
400 (1)
Unloading the ball
200 Start sliding
100
(2)
0 0
1000
(b) 800 Current density (μA cm-2)
Current density (μA cm-2)
300
End sliding
2000 3000 Time (s)
4000
600 500
(1)
Start sliding
400
Unloading the ball
(2)
300 200 0
1000
2000 3000 Time (s)
4000
5000
Start sliding
-20 Unloading the ball
(2)
-40 -60
End sliding
(1)
-80 0
5000
End sliding
700
0
(b)
0
Current density (μA cm-2)
Current density (μA cm-2)
(a)
-20
1000
2000 3000 Time (s)
Start sliding
-40
4000
5000
Unloading the ball
(2)
-60 -80
(1)
End sliding
-100 0
1000
2000 3000 Time (s)
4000
5000
Fig. 8. Evolution of current density readings of (a) amorphous and (b) nanocrystalline CoeW coatings under anodic polarisation (E ¼ þ0.1 V vs. OCP) before, after and during tribocorrosion experiments (1) with and (2) without touching the specimen surface in 3.5 wt-% NaCl solution.
Fig. 9. Evolution of current density readings of (a) amorphous and (b) nanocrystalline CoeW coatings under cathodic polarisation (E ¼ 0.7 V/Ag/AgCl) before, after and during tribocorrosion experiments (1) with and (2) without touching the specimen surface in 3.5 wt-% NaCl solution.
oxidation process occurs more quickly which encourages the growth of passive film [31]. Subsequently, thickening of the passive film causes a decrease in the conductivity of oxide layer, an increase in the band gap, and thus, a significant decrease in the rate of oxidation process. This stabilizes the current density and it remains constant until the ball motion is started. By the ball motion, a sharp increase in current density was observed due to mass transport effect in non-contact condition and also depassivation in sliding with contact condition. However, the current density decreased rapidly to its initial values when the ball motion finished. By unloading the ball from the specimen surface, a slight increase in current density followed by a gradual decrease was observed for both amorphous (Fig. 8a) and nanocrystalline (Fig. 8b) coatings. Under anodic polarisation, the decrease in passive film thickness may be responsible for appearance of a peak at the start of ball motion in non-contact condition (Fig. 8a and b). Over the course of sliding, owing to the gradual decrease of OCP for the amorphous coating (Fig. 7a), it should experience higher anodic polarisation, and thus, thickening and restoration of passive film is expected. The minor alteration in anodic current of the amorphous coating under anodic polarisation (Fig. 8a) indicates that the passive film formed on the surface of the amorphous coating is more protective than that formed on the nanocrystalline one. Thus, the effect of sliding (or ball motion alone) on electrochemical degradation is less intense for the amorphous coating, because lower anodic currents were achieved (Fig. 8a and b). Fig. 9 shows current density readings under cathodic polarisation. Under cathodic polarisation, the passive film can be totally (or partially) reduced from the coatings surface. Moreover, no significant corrosion attack is anticipated because of the cathodic protection of specimen. This fact could be observed in Fig. 9, where
the current density drastically approaches very low values by applying the cathodic potential. By the sliding, the current density became more cathodic probably due to the acceleration of cathodic reactions such as hydrogen evolution on the wear scars and also mass transport effect. At the end of sliding, the current density returned to its initial magnitude in both contact conditions. Again, by unloading the ball from specimen surface, an increase in current density followed by a gradual decrease is observed for both amorphous (Fig. 9a) and nanocrystalline (Fig. 9b) coatings. In this case, the ball motion eliminates the concentration polarisation, and hence the rate of the hydrogen evolution reaction (or the reduction of water) goes up. Because of the absence of considerable corrosion reaction at cathodic polarisation, volume loss of the coatings is mainly mechanical wear and proportion of corrosion accelerated wear (CAW) is negligible. Indeed, there is an inverse relation between the wear rate and hardness of the coatings at cathodic condition [30]. During the sliding, the recorded anodic current density, ia, in Fig. 8 could be achieved by: ia ¼ iw þ ip
(3)
where iw is the current density arises from the ball motion and ip is the anodic current density in the absence of sliding. Furthermore, iw can be considered as: iw ¼ iMT þ iWAC
(4)
where iMT and iWAC are the current densities originated from mass transport effect (MT) and wear accelerated corrosion (WAC), respectively. Table 3 summarizes the proportion of WAC to MT for
Please cite this article in press as: N. Fathollahzade, K. Raeissi, Electrochemical evaluation of corrosion and tribocorrosion behaviour of amorphous and nanocrystalline cobaltetungsten electrodeposited coatings, Materials Chemistry and Physics (2014), http://dx.doi.org/ 10.1016/j.matchemphys.2014.07.013
N. Fathollahzade, K. Raeissi / Materials Chemistry and Physics xxx (2014) 1e10 Table 3 The ratio of wear accelerated corrosion (WAC) to mass transport (MT) phenomenon in evolution of tribocorrosion data plotted in Figs. 7e9. Coating structure
Conditions OCP
Amorphous Nanocrystalline
Z
Cathodic polarisation
WAC
MT
WAC
MT
WAC
MT
9% 4%
91% 96%
22% 17%
78% 83%
18% 14%
82% 86%
Z ia dt ¼
Z iMT dt þ
qa ¼ qMT þ qWAC þ qp
provides stronger hydrogen evolution probably due to the induced stresses of cold-working [32]. 3.5. Characterisation of wear scars at OCP conditions
Anodic polarisation
the coatings in sliding at OCP, anodic and cathodic conditions. These values were calculated using the integration of current density vs. time at sliding time intervals. For example in the case of anodic polarisation (curve (1) in Fig. 8), the total charge, qa, is given by:
qa ¼
7
Z iWAC dt þ
ip dt
(5)
(6)
In other words, qa is calculated by integration of curve (1) in Fig. 8. The total charge in no contact condition, qb, is the sum of mass transport charge, qMT, and the anodic charge in sliding condition, qp, and is calculated by integration of curve (2) in Fig. 8. However, the anodic charge, qp, is found from Faraday's low. On the other hand, the charge of WAC, qWAC, equals qa qb. At OCP condition, the ratios in Table 3 were determined by averaging the OCP values during the sliding. The higher proportion of mass transport in tribocorrosion experiments (Table 3) implies that the dissolution reaction of the coatings is mainly a mass-transport controlled process. In all conditions, the proportion of WAC of the amorphous coating is greater than that of the nanocrystalline one. A possible reason is lower hardness of the amorphous coating compared to the nanocrystalline one, which is responsible for its higher damage. Under anodic polarisation, a higher damage is revealed by a wider wear scar. On the other hand, under cathodic polarisation, a higher damage
Line profiles and SEM observations of wear scars after the tribocorrosion experiments give useful data in order to determine the shape and width of the wear scars for calculating total volume loss of the working electrode. Figs. 10 and 11 respectively show SEM images and line profiles of the wear scars created at OCP condition for the amorphous and nanocrystalline CoeW coatings. The width and depth of the wear scars in the amorphous coating were estimated to be about 390 and 4.5 mm, respectively, whilst they were 340 and 3.3 mm for the nanocrystalline coatings. Considering the shape of wear scar as a part of a cylinder, the total volume loss of the amorphous and nanocrystalline coatings was estimated to be around 6 103 and 4 103 mm3, respectively. The higher total volume loss of the amorphous coating is in agreement with its lower microhardness compared to the nanocrystalline coating. The lower microhardness of the amorphous coating results higher percentage of WAC compared to the nanocrystalline one. After the tribocorrosion experiment, some microcracks were detected in the vicinity of the wear scars even in the nanocrystalline coating (Fig. 10b), where no microcracks were observed before sliding (Fig. 4b). The formation of these superficial microcracks in the vicinity of wear scars might be another reason for the difference in OCP values observed at the end of tribocorrosion experiments comparing its initial values (Fig. 7). These microcracks were extended or created by applying the load due to the coatings brittleness. According to SEM images of the coatings (Fig. 10), it seems that the generated microcracks are superficial. Thus, for both coatings, a surface fatigue wear is associated (Fig. 10). According to Fig. 10, the aggregation of wear debris in environs of the wear scars in the amorphous coating is more than that of the nanocrystalline one. 3.6. EIS measurements Fig. 12 shows Nyquist and Bode plots of the amorphous and nanocrystalline CoeW coatings before and after the tribocorrosion
Fig. 10. SEM images (BSE) of wear scars of (a) amorphous and (b) nanocrystalline coatings after the tribocorrosion experiments at OCP condition.
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Fig. 11. Line profiles of wear scars of amorphous (A) and nanocrystalline (N) CoeW coatings (Reconstructed Figure).
experiments at OCP. The appearance of two time constants could be easily distinguished from the plots. If the coatings have pores and cracks, they can also interfere in corrosion process lead to a capacitive response at low frequencies [33]. In Nyquist plots (Fig. 12), the high-frequency loop (the smaller one) is the response of passive film/electrolyte interface, whilst the low-frequency loop (the larger one) is connected to the coating defects and arises due to a finite-length diffusion process of reactants inside the coating defects [34]. The electrical elements extracted from Nyquist plots using the equivalent circuit model depicted in Fig. 13 are listed in Table 4. In the equivalent circuit, R(s) denotes the uncompensated solution resistance which depends upon the distance between reference and working electrodes and solution conductivity. The potentials of points A, Y and B (Fig. 13) represent the potentials at the reference electrode, solution side of double layer and coating surface,
respectively. In order to simulate the effect of passive layer at high frequencies, R(ox) and CPE(ox) were applied which correspond to the oxide resistance and the constant phase element (CPE), respectively. In order to represent the deviation of a capacitive loop from the ideal capacitive behaviour (i.e. depression), it is common to replace the capacitors with CPEs. According to the literature [35], this deviation has been related to the heterogeneities of surface or non-uniform current distribution. The higher R(ox) of the passive film formed on the amorphous coating at OCP (Table 4) confirms the presence of a less defective passive film. On the other hand, R(d) and CPE(d) are equivalent to the diffusion resistance and CPE of diffusion process, which are connected to the low-frequency loop [34]. As is seen in Table 4, a smaller value of CPE(d) Power (0.7595) is obtained for the amorphous coating before sliding. Hence, the centre of its lowfrequency loop fell lower than that for the nanocrystalline coating (Fig. 12a). This more probably arises due to the presence of the microcrack network in the amorphous coating affecting the diffusion process. EIS experiments also provide a complete study on the role of intermediate species adsorbed on a corroding surface [36]. Typically, the magnitude of CPE(ox) should be in order of microFarads per square centimetre [37], but as is seen in Table 4, it has increased significantly to a higher order, particularly before tribocorrosion tests. The main reason is most likely the participation of the adsorbed intermediates in the interfacial corrosion process [33] (e.g. Co(OH)ads in cobalt dissolution process). However, the order of CPE(d) is milliFarads per square centimetre as supposed for finite-length diffusion processes [37]. The polarization resistance, R(p), is the sum of resistances and defined as the impedance limit at zero frequency found by replacing the capacitors (or CPEs) with open circuits [38]. Table 4 reveals that before and after tribocorrosion experiments, R(p) values of the amorphous coating are higher than those of the nanocrystalline coating. As R(p) is inversely related to the corrosion rate (Eq. (2)), the corrosion resistance of the amorphous coating is higher than that of the nanocrystalline one. This is in accordance to
(a) 500
90 Fit result (A)
Phase angle (deg)
-Zimg (Ω cm2)
400
(N)
Fit result (N)
300
(A)
200 100 0
0
(A)
50 30
(N)
10 -10 -30
20 40 60 80 100 120 140 160 180
0.1
Zre (Ω cm2)
(b) 700
10
1000
100000
Frequency (Hz) 70
600
Fit result (A)
500
Fit result (N)
Phase angle (deg)
-Zimg (Ω cm2)
70
400 300 200 100 0 0
150
300
450
600
Zre (Ω cm2)
750
900
50 30 10 -10 -30 0.1
10
1000
100000
Frequency (Hz)
Fig. 12. Nyquist (on the left) and Bode (on the right) plots of the amorphous (A) and nanocrystalline (N) CoeW coatings (a) before and (b) after tribocorrosion experiments at OCP condition in 3.5 wt-% NaCl solution.
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N. Fathollahzade, K. Raeissi / Materials Chemistry and Physics xxx (2014) 1e10
Fig. 13. The equivalent circuit applied for EIS plots in Fig. 12. (R: resistor, CPE: constant phase element).
the results obtained by potentiodynamic polarisation readings and the values of R(p) obtained from EIS plots are very similar to those found from polarization curves (Table 2). According to Table 4, R(ox) of both coatings has decreased after sliding. As R(ox) stems from the entire surface of test specimen (worn and unworn areas), the decrease of R(ox) means that the passive layer formed in worn area after the end of sliding, has not yet regained its initial thickness and/or protective properties [39]. 4. Conclusions In this investigation, the amorphous and nanocrystalline CoeW coatings with about 15 mm thickness were electrodeposited on copper substrate. The tungsten content of the nanocrystalline coating increased from ~35 to ~48 wt-% by increasing the deposition current density and bath temperature where the amorphous coating was obtained. The grain size of the nanocrystalline CoeW coating was estimated to be about 15 nm. A nodular morphology was exhibited for both amorphous and nanocrystalline CoeW coatings, but a microcrack network was only observed in the amorphous CoeW deposit. Briefly, it could be concluded that: 1) The amorphous coating was more resistant to corrosion attacks in 3.5 wt-% NaCl solution compared to the nanocrystalline one and the microcracks plugging could prevent the build-up of galvanic corrosion cells, and thus eliminate the negative effect of the microcracks on corrosion resistance of the amorphous coating. 2) During OCP sliding, the potential of amorphous coating gradually became more active probably due to the widening of wear scar, and thus expansion of active area. This is attributed to the lower microhardness of amorphous coating compared with the nanocrystalline one. 3) Tribocorrosion behaviour of CoeW coatings at OCP could be explained in terms of mass transport (MT) and wear accelerated corrosion (WAC). The proportion of MT in tribocorrosion was higher than WAC, which implied that the corrosion reaction of the coatings was mainly a mass transport controlled process.
Table 4 Data extracted from EIS plots in Fig. 12 using the equivalent circuit in Fig. 13. Element
Value Before tribocorrosion
After tribocorrosion
Amorphous Nanocrystalline Amorphous Nanocrystalline R(s) (U cm2) CPE(ox) T (F cm2) CPE(ox) Power R(ox) (U cm2) CPE(d) T (F cm2) CPE(d) Power R(d) (U cm2) R(p) (U cm2)
1.57 0.002388
1.88 0.005291
4.83 0.001206
12.74 0.001254
0.6435 40.18 0.02289
0.5479 19.43 0.02477
0.3622 7.28 0.001155
1.0000 8.39 0.002893
0.7595 2940 2980
0.9000 2090 2110
0.8661 2447 2454
0.6095 1835 1843
9
3) In anodic sliding, the evolution of current readings indicates that the effect of sliding (or ball motion alone) on electrochemical degradation is more intense for the nanocrystalline coating, because of the higher achieved anodic currents. In this case, the passive film formed on the surface of the amorphous coating was more protective than that formed on the nanocrystalline one. However, the volume loss of the amorphous coating was higher than that of the nanocrystalline one. This is related to the higher microhardness of nanocrystalline coating. 4) Under cathodic polarisation, the passive film could be totally (or partially) reduced in tribocorrosion experiment, but no remarkable corrosion was observed. In this case, the ball motion eliminated the concentration polarisation, and hence the rate of the hydrogen evolution reaction (or the reduction of water) went up. References [1] Z. Ghaferi, K. Raeissi, M.A. Golozar, H. Edris, Characterization of nanocrystalline CoeW coatings on Cu substrate, electrodeposited from a citrate-ammonia bath, Surf. Coat. Technol. 206 (2011) 497e505. [2] P. Lammel, L.D. Rafailovic, M. Kolb, K. Pohl, A.H. Whitehead, G. Grundmeier, B. Gollas, Analysis of rain erosion resistance of electroplated nickeletungsten alloy coatings, Surf. Coat. Technol. 206 (2012) 2545e2551. [3] L. Wang, Y. Gao, T. Xu, Q. Xue, A comparative study on the tribological behaviour of nanocrystalline nickel and cobalt coatings correlated with grain size and phase structure, Mater. Chem. Phys. 99 (2006) 96e103. [4] F. Su, C. Liu, P. Huang, Establishing relationships between electrodeposition techniques, microstructure and properties of nanocrystalline CoeW alloy coatings, J. Alloys. Compd. 557 (2013) 228e238. [5] F. Su, P. Huang, Microstructure and tribological property of nanocrystalline CoeW alloy coating produced by dual-pulse electrodeposition, Mater. Chem. Phys. 134 (2012) 350e359. [6] D.P. Weston, S.J. Harris, H. Capel, N. Ahmed, P.H. Shipway, J.M. Yellup, Nanostructured CoeW coatings produced by electrodeposition to replace hard Cr on aerospace components, Trans. Inst. Met. Finish. 88 (2010) 47e56. [7] H. Capel, P.H. Shipway, S.J. Harris, Sliding wear behaviour of electrodeposited cobaltetungsten and cobaltetungsteneiron alloys, Wear 255 (2003) 917e923. [8] D.P. Weston, P.H. Shipway, S.J. Harris, M.K. Cheng, Friction and sliding wear behaviour of electrodeposited cobalt and cobaltetungsten alloy coatings for replacement of electrodeposited chromium, Wear 267 (2009) 934e943. [9] F. Su, C. Liu, P. Huang, Friction and wear of nanocrystalline Co and CoeW alloy coatings produced by pulse reverse electrodeposition, Wear 300 (2013) 114e125. [10] A. Igual Munoz, N. Espallargas, Tribocorrosion mechanisms in sliding contacts, in: D. Landolt, S. Mischler (Eds.), Tribocorrosion of Passive Metals and Coatings, Woodhead Publishing Limited, 2011, p. 118. [11] P. Jemmely, S. Mischler, D. Landolt, Electrochemical modeling of passivation phenomena in tribocorrosion, Wear 237 (2000) 63e76. [12] E. Arslan, Y. Totik, I. Efeoglu, The investigation of the tribocorrosion properties of DLC coatings deposited on Ti6Al4V alloys by CFUBMS, Prog. Org. Coat. 74 (2012) 768e771. [13] N. Sato, T. Ohtsuka, Anodic oxidation of cobalt in neutral and basic solution, J. Electrochem. Soc. 125 (1978) 1735e1740. [14] M. Metikos-Hukovic, R. Babic, Passivation and corrosion behaviours of cobalt and cobaltechromiumemolybdenum alloy, Corros. Sci. 49 (2007) 3570e3579. [15] M. Pontinha, S. Faty, M.G. Walls, M.S. Ferreira, M. Da Cunha Belo, Electronic structure of anodic oxide films formed on cobalt by cyclic voltammetry, Corros. Sci. 48 (2006) 2971e2986. [16] H. Jung, A. Alfantazi, Corrosion properties of electrodeposited cobalt in sulfate solutions containing chloride ions, Electrochim. Acta 55 (2010) 865e869. [17] A.M. Human, B. Roebuck, H.E. Exner, Electrochemical polarisation and corrosion behaviour of cobalt and Co(W,C) alloys in 1 N sulphuric acid, Mater. Sci. Eng. A 241 (1998) 202e210. [18] A. Bodaghi, J. Hosseini, Corrosion behavior of electrodeposited cobaltetungsten alloy coatings in NaCl aqueous solution, Int. J. Electrochem. Sci. 7 (2012) 2584e2595. [19] Z. Abdel Hamid, Electrodeposition of cobaltetungsten alloys from acidic bath containing cationic surfactants, Mater. Lett. 57 (2003) 2558e2564. [20] R.C. Santana, A.N. Campos, E.A. Medeiros, A.M. Oliveira, L.F. Oliveira, S. Prasad, Studies on electrodeposition and corrosion behaviour of a NieWeCo amorphous alloy, J. Mater. Sci. 42 (2007) 9137e9144. [21] H.J. Yun, S. Dulal, C.B. Shin, C. Kim, Characterisation of electrodeposited CoeWeP amorphous coatings on carbon steel, Electrochim. Acta 54 (2008) 370e375. [22] S.K. Ghosh, J.P. Celis, Tribological and tribocorrosion behaviour of electrodeposited CoW alloys and CoWeWC nanocomposites, Tribol. Int. 68 (2013) 11e16.
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[23] Y. Liu, S. Shrestha, W.E. Mustain, Synthesis of nanosize tungsten oxide and its evaluation as an electrocatalyst support for oxygen reduction in acid media, ACS Catal. 2 (2012) 456e463. [24] S. Hassani, K. Raeissi, M. Azzi, D. Li, M.A. Golozar, J.A. Szpunar, Improving the corrosion and tribocorrosion resistance of NieCo nanocrystalline coatings in NaOH solution, Corros. Sci. 51 (2009) 2371e2379. [25] R.C. Santana, S. Prasad, A.N. Campos, F.O. Araujo, G.P. Da Silva, P. De LimaNeto, Electrodeposition and corrosion behaviour of a NieWeB amorphous alloy, J. Appl. Electrochem. 36 (2006) 105e113. [26] A. Singh, N. Iyer, S.P. Harimkar, Prediction of glass forming ability and glass forming range for electrodeposited binary CoeW alloys, J. Appl. Phys. 111 (2012) 103508. [27] R.W. Revie, H.H. Uhlig, Kinetics: Polarization and Corrosion Rates, Corrosion and Corrosion Control: an Introduction to Corrosion Science and Engineering, John Wiley & Sons, New Jersey, 2008, p. 72. [28] J. Stojadinovic, D. Bouvet, M. Declercq, S. Mischler, Effect of electrode potential on the tribocorrosion of tungsten, Tribol. Int. 42 (2009) 575e583. [29] M.-G. Verge, C.-O.A. Olsson, D. Landolt, Anodic oxide growth on tungsten studied by EQCM, EIS and AES, Corros. Sci. 46 (2004) 2583e2600. [30] D. Landolt, S. Mischler, M. Stemp, Electrochemical methods in tribocorrosion: a critical appraisal, Electrochim. Acta 46 (2001) 3913e3929. [31] C.M.A. Brett, A.M.O. Brett, The Semiconductor Electrode: the Space-charge Region, Electrochemistry Principles, Methods, and Applications, Oxford University Press, 1993, pp. 58e64.
[32] L.L. Shreir, R.A. Jarman, G.T. Burstein, Dislocations, Etch Pits and the Effect of Cold Work on Corrosion, Corrosion e Metal/Environment Reactions, vol. 1, Butterworth-Heinemann, 1994, pp. 37e39. [33] R.G. Kelly, J.R. Scully, D.W. Shoesmith, R.G. Buchheit, Influence of Corrosion Product Deposits, Electrochemical Techniques in Corrosion Science and Engineering, Marcel Dekker, New York, 2002, pp. 223e230. [34] C. Liu, Q. Bi, A. Leyland, A. Matthews, An electrochemical impedance spectroscopy study of the corrosion behaviour of PVD coated steels in 0.5 N NaCl aqueous solution: part I. Establishment of equivalent circuits for EIS data modelling, Corros. Sci. 45 (2003) 1243e1256. [35] N. Mary, C. Alemany-Dumont, B. Normand, O. Heintz, V. Broudic, C. Jegou, Electrochemical behaviour of copperenickel alloys as immobilisation matrices for the storage of fission products in CO2-enriched synthetic groundwater, Electrochim. Acta 98 (2013) 11e19. [36] P. Ponthiaux, F. Wenger, D. Drees, J.P. Celis, Electrochemical techniques for studying tribocorrosion processes, Wear 256 (2004) 459e468. [37] W.S. Tait, Analyzing and Interpreting EIS Spectra, an Introduction to Electrochemical Corrosion Testing for Practicing Engineers and Scientists, PairODocs Publication, 1994, pp. 95e114. [38] D.A. Harrington, P. Van Den Driessche, Mechanism and equivalent circuits in electrochemical impedance spectroscopy, Electrochim. Acta 56 (2011) 8005e8013. [39] M. Azzi, M. Paquette, J.A. Szpunar, J.E. Klemberg-Sapieha, L. Martinu, Tribocorrosion behaviour of DLC-coated 316L stainless steel, Wear 267 (2009) 860e866.
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Electrochemical evaluation of corrosion and tribocorrosion behaviour of amorphous and nanocrystalline cobaltetungsten electrodeposited coatings N. Fathollahzade, K. Raeissi* Department of Materials Engineering, Isfahan University of Technology, Isfahan 84156-83111, Iran
h i g h l i g h t s Mass-transport effect had higher proportion in tribocorrosion of CoeW coatings. The major electrochemical-wear degradation was for the nanocrystalline coating. The higher proportion of wear accelerated corrosion was for the amorphous coating. Superficial microcracks were formed near scars due to the coatings brittleness.
a r t i c l e i n f o
a b s t r a c t
Article history: Received 4 November 2013 Received in revised form 4 May 2014 Accepted 9 July 2014 Available online xxx
Amorphous and nanocrystalline CoeW coatings were electrodeposited on copper substrates from a citrateeammonia bath. The coatings showed nodular surface morphologies, but a microcrack network was detected in the amorphous coating. However, a better corrosion resistance was achieved for the amorphous coating. During sliding under open circuit potential (OCP) condition, the potential of amorphous coating gradually became more active probably due to the widening of wear scar, and thus expansion of active area. The amorphous coatings showed a higher volume loss at OCP probably due to its lower microhardness. In anodic sliding, a sharp increase in current density was observed due to mass transport and depassivation effects. In all sliding conditions, the proportion of mass transport was higher than wear accelerated corrosion, which implied that the dissolution reaction of the coatings was mainly a mass-transport controlled process. The results also showed that the effect of sliding on degradation is more intense for the nanocrystalline coating. For both coatings, the formation of the superficial microcracks in the vicinity of wear scars indicating on a surface fatigue wear mechanism. © 2014 Elsevier B.V. All rights reserved.
Keywords: Coatings Amorphous materials Nanostructures Corrosion Tribology
1. Introduction Cobalt and nickel based alloys containing tungsten have the potential to become promising replacements for hard chromium coatings deposited from toxic, carcinogenic and environmentally unfriendly hexavalent chromium baths [1,2]. Wang et al. [3] reported improved tribological properties of electrodeposited Co over Ni coatings with hexagonal close packed (hcp) and face centred cubic (fcc) structures, respectively, and attributed this result to the higher resistance of hcp structure to adhesion interactions during the wear process. Higher W content in bath, smaller grain size [4]
* Corresponding author. Fax: þ98 3113912752. E-mail address: [email protected] (K. Raeissi).
and strong hcp (100) crystal orientation [5] improve the hardness of the nanocrystalline CoeW coatings which is higher than that of the amorphous structures [6]. Su et al. [5] obtained a microhardness range of 800e1000 HV for nanocrystalline CoeW alloys electrodeposited under various conditions, while CoeW coatings produced by Ghaferi et al. [1] had a microhardness of 280e620 HV. Capel et al. [7] reported an increase in hardness (from 540 to 1200 HV) and a significant decrease in corrosion resistance by heat treatment. Some authors found a lower rate of wear for CoeW electrodeposit compared to pure cobalt electrodeposit [8] on account of the incorporations of W, and thus the increased hardness [9]. Tribocorrosion is the degradation of materials due to the simultaneous effects of electrochemical corrosion and mechanical wear in a tribological contact [10]. The total material removal rate
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N. Fathollahzade, K. Raeissi / Materials Chemistry and Physics xxx (2014) 1e10
differs from sum of the wear rate measured in the absence of corrosion and the corrosion rate observed in absence of wear [11]. For example, passive coatings would subject to active corrosion due to the annihilation of protective passive film in wet sliding contact, when the frictional movement in corrosive medium is continuous [12]. Regardless of the considerable corrosion data on cobalt and its alloys available in the literature [1,13e21], the tribocorrosion behaviour of them have been less investigated. Recently, Ghosh and Celis [22] reported that the open circuit potential (OCP) of CoeW alloy in 0.5 M NaCl electrolyte became nobler during the sliding and remained at that nobler value even after the end of sliding test. However, the OCP decreased progressively and reached the value recorded before starting the sliding test until the corundum ball was unloaded and removed from the sample surface [22]. These researchers [22] attributed this behaviour to the main role of tungsten in CoeW coatings which forms WO3 and catalyses the oxygen reduction reaction (ORR). Nevertheless, Liu et al. [23] suggested WO3 as a possible support material for monolayer Pt ORR electrocatalysts in acid electrolytes and have no idea about catalytic effect of WO3 on ORR. In this study, corrosion and tribocorrosion behaviours of amorphous and nanocrystalline CoeW coatings electrodeposited from a citrate-ammonia bath are investigated. The evolution of OCP values, anodic and cathodic current readings and SEM images taken from wear scars are applied to determine tribocorrosion behaviour of the coatings. The tribocorrosion resistance of the coatings were also estimated by calculating the volume loss in worn track using line profilometry.
Fig. 1. Schematic representation of reciprocating ball-on-plate tribometre. (1) DC motor, (2) specimen (working electrode), (3) counter electrode, (4) reference electrode, (5) load cells, (6) ball and its holder, (7) spring, (8) loading bolt (9) connecting rod, (10) potentiostat.
electromagnetic (eddy-current) method. The grain size of the coatings was calculated from X-ray diffraction (XRD) patterns obtained by an X-ray diffractometer (model X'pert Philips) with Cu Ka radiation (l ¼ 1.5418 Å), using the Scherrer equation. For the curve fitting analysis required for Full-Width at Half Maximum (FWHM) determination, the four-parameter Gaussian function was employed [24]. The instrumental line broadening was measured using a silicon peak corrected by following equation:
b/bexp ¼ 1 (bins/bexp)2
where b is the modified FWHM, bexp is the calculated FWHM of experiment specimen and bins is the FWHM of monocrystalline silicon. A Philips CM120 transmission electron microscope (TEM) with 0.037 Å wavelength and 100 kV voltage was used to detect the grain size. TEM sample preparation was performed by scratching the coating from substrate using a surgical blade and the resulted powders were dispersed in alcohol to choose a suitable particle for analysis.
(101)
Cu
(b)
Table 1 Composition of the bath for CoeW electrodeposition. Substance
Amount (g l1)
CoSO4$7H2O Na2WO4$2H2O NH4CL NaBr Na3C6H5O7$2H2O
40.24 51.22 26.74 14.34 178.58
Cu
(002)
The copper used as substrate was a disk of 0.9 cm2 surface area. The substrate surface was abraded with wet SiC papers up to grit size no. 1200 and then anodically polished for 15 min in a bath containing 80 vol.-% concentrated phosphoric acid and 20 vol.-% distilled water at 1.4 V, 25 C and 700 rpm. The cathode was a large plate of stainless steel. Immediately after washing with distilled water, the specimen (as the working electrode) was introduced into a cell containing electrodeposition solution with the composition shown in Table 1. The pH of the bath was adjusted at 7.5 using dilute NaOH and H2SO4 solutions. A platinum plate was applied as a counter electrode. The electrodeposition process was performed using a digital coulometer (BHP 2050). The amorphous and nanocrystalline CoeW alloys were electrodeposited at deposition current densities of 70 and 15 mA cm2 and temperatures of 45 and 25 C, respectively. The amorphous coatings were electrodeposited at 110 rpm, whereas nanocrystalline coatings were obtained at stationary conditions. The times for amorphous and nanocrystalline CoeW electrodeposition were 500 and 2330 s, respectively. Philips XL30 scanning electron microscope (SEM) was used to study the surface morphology and topography of the coatings before and after tribocorrosion experiments. Elemental compositions of the coatings were determined using an electron dispersive spectroscope (EDS) coupled with a Seron Technology AIS 2300C. Thickness of the coatings was measured about 15 mm using the
(1)
(100)
2. Materials and experimental methods
Intensity
2
(a) 30
35
40
45
50
55
60
2θ (degree) Fig. 2. XRD patterns of electrodeposited (a) amorphous (with 48 wt-% W) and (b) nanocrystalline (with 35 wt-% W) CoeW coatings.
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N. Fathollahzade, K. Raeissi / Materials Chemistry and Physics xxx (2014) 1e10
3
were recorded using an EG&G potentiostat/galvanostat (model 263A). Tribocorrosion tests under anodic and cathodic polarisation were done by applying constant potentials of þ0.1 V/OCP and 0.7 V/Ag/AgCl, respectively. The specimen was relaxed for 1200 and 1000 s before and after sliding, respectively, then the ball was unloaded and the recording were continued further for 1000 s. The speed and stock length of a complete cycle of reciprocating movement on the specimen surface were 1 cm s1 and 1 cm, respectively. After sliding, line profiles of wear scars were drawn using a surface profilometer (Mitutoyo Surftest SJ-301), in order to determine the shape and depth of the wear scars, and calculate the volume loss of specimen. In order for simulation of the fluid dynamics as in real tribocorrosion experiments, the ball was placed at a distance of about 0.5 mm above the surface of the working electrode and ball motion was done without any contact with the specimen surface under OCP and anodic/cathodic polarisation conditions. Each measurement was repeated for three times. 3. Results and discussion Fig. 3. Bright field of TEM image from nanocrystalline coating.
Microhardness of the coatings was determined using Buehler Micromet 5101 by applying a constant load of 25 g for 10 s. The calculated value of microhardness was the average number of five random points on the surface of each coating. Each measurement was repeated for three times. All electrochemical experiments were done using platinum as the auxiliary (counter) electrode, saturated Ag/AgCl as the reference electrode (with offset potential of 0.197 V/SHE) and the coated specimen as the working electrode. The testing electrolyte was 3.5 wt-% NaCl solution applied at room temperature in stationary condition. The tests were run after 60 min immersion of the working electrode. In all representations, cathodic currents have been defined as negative values. Five specimens were taken for each electrochemical experiment. Potentiodynamic polarisation measurements were carried out using a Princeton Applied Research (PARSTAT) potentiostat model 2273 at a sweep rate of 1 mV s1. Electrochemical impedance spectroscopy (EIS) measurements were run at OCP, before and after sliding, using an EG&G AC responser (model 1025) coupled with an EG&G potentiostat/galvanostat (model 263A). The voltage amplitude and the frequency range were 10 mV and 100 kHze10 mHz, respectively. Tribocorrosion experiments were performed using a reciprocating ball-on-plate tribometre described in Fig. 1. The working electrode surface was rubbed with a SiC ball under the load of 2 N for 1800 s at 1 Hz. The OCP values and measured current readings
3.1. Characterisation of CoeW coatings XRD patterns of the electrodeposited amorphous and nanocrystalline CoeW coatings are shown in Fig. 2. The pattern of nanocrystalline coating shows three characteristic peaks around 2q of 41.3, 44.3 and 47 which are very close to those in JCPDS 05-0727 attributed to (100), (002) and (101) planes of hcp Co, respectively. However, the shifts in peak positions of hcp Co to lower 2q is observed in Fig. 2 and related to the formation of a supersaturated solid solution of W in Co. In this case, since Co atoms in the lattice are substituted by bigger W atoms, an increase in the average closest distance between two adjacent atoms should be expected. On the other hand, a broad peak around 2q range of 40e50 is observed in the pattern of amorphous coating that confirms the amorphous structure of the alloy. Since, the characteristic peaks of W and Co3W are absent in Fig. 2, thus complete dissolution of tungsten in cobalt has occurred. For both coatings, the characteristic peaks of copper substrate are evidenced around 2q of 43.3 and 50.4 belonging to (111) and (200) planes of fcc copper which are consistent with cubic Cu lines in PDF 04-0836. Using the Full-Width at Half Maximum (FWHM) of Co (100) peak and the Sherrer equation, the grain size of nanocrystalline CoeW coating was estimated to be around 15 nm. The TEM result also confirms a same average grain size as is seen in Fig. 3. A nodular surface morphology can be observed in SEM images of the amorphous and nanocrystalline CoeW electrodeposited coatings (Fig. 4). Furthermore, a microcrack network is evidenced in the amorphous CoeW coating. The formation of the microcrack network could be related to the extensive hydrogen evolution
Fig. 4. SEM images of the area used for EDS analysis of (a) amorphous and (b) nanocrystalline CoeW coating. The higher magnification pics are also inserted.
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4
N. Fathollahzade, K. Raeissi / Materials Chemistry and Physics xxx (2014) 1e10 A - BNC
A - NS
A - SC
N - SC
N - NS
S - NS
stationary and ball motion conditions are shown in Fig. 5. Data extracted from Fig. 5 using Tafel extrapolation method are listed in Table 2. The polarisation resistance, R(p), is calculated using the SterneGeary equation [27]:
N - BNC
0.01
Current density (A cm-2)
0.001
R(p) ¼ (babc)/(2.3ic(ba þ bc))
0.0001
0.00001
0.000001
0.0000001 -0.9
-0.8
-0.7
-0.6
-0.5
-0.4
-0.3
-0.2
-0.1
0
Potential (V/Ag/AgCl) Fig. 5. Potentiodynamic polarisation curves (E vs. log i) of Cu substrate (S) and the amorphous (A) and nanocrystalline (N) CoeW coatings under different sliding condition in 3.5 wt-% NaCl solution. (NS: no sliding, SC: sliding with contact, BNC: ball motion with no contact).
coming up by increasing the current density and temperature of deposition process [1]. However, some authors have attributed this to the internal stresses built up by the presence of a high percentage of tungsten in the amorphous coating [25]. EDS analysis revealed that the amorphous and nanocrystalline CoeW electrodeposits had about 48 and 35 wt-% tungsten, respectively. It is reported that the transition from nanocrystalline phase to amorphous depends on tungsten content and can occur when it exceeds a critical value (about 45 wt-%) making the disordered state energetically more favourable [26]. Despite of the higher content of tungsten in amorphous coating, nanocrystalline coating had higher microhardness value. Vickers microhardness of the amorphous and nanocrystalline CoeW coatings were about 290 and 360 HV, respectively, which are lower than those reported by some authors [5,7,9]. It has been reported that the hardness of CoeW alloy coatings greatly depends on the content of W, the grain size, crystal orientation and phase structure [5,9] which are influenced by electrodeposition conditions such as bath composition and deposition current density. Incorporations of W in Co up to ~20 wt-% [5] causes a significant increase in the hardness of CoeW alloys owing to the decrease of the grain size, however, larger incorporation of W leads to a less strong hcp (100) orientation [5] which counteracts with the decrease of the grain size and thus decreases the hardness. 3.2. Potentiodynamic polarisation measurements Potentiodynamic polarisation plots (E vs. log i) for Cu substrate, the amorphous and nanocrystalline CoeW electrodeposits at
(2)
where ba and bc are Tafel constants for the anodic and cathodic reactions, respectively, and ic is the corrosion current density. As is seen in Table 2, the coatings which are under ball motion (with or without contact with the specimen surface) show lower R(p) with E(corr) values shifted to the more active values. This is more intense in contact (sliding) condition. However, the changes observed in OCP values were negligible (Table 2). According to R(p) values, the amorphous CoeW coating showed better corrosion resistance than the nanocrystalline one. In each case, E(corr) of the amorphous coating was nobler than that of the nanocrystalline one, whereas the nanocrystalline coating showed nobler OCP (Table 2). The main reason behind the difference exists between OCP and E(corr) of the coatings (Table 2) is more likely the reduction of passive layer whilst drawing the cathodic branch in potentiodynamic polarisation curves, which causes a shift in mixed potential (E(corr)) to a more negative (active) value. In the absence (or weakness) of passive layer on the coatings surface, the amorphous coating should act nobler than the nanocrystalline one at least due to the presence of microcracks, and hence interference of the nobler Cu substrate in measurements. In the presence of passive layer, which may plug most of the microcracks in amorphous coating (Fig. 6a), the influence of Cu substrate is diminished in OCP measurements. Fig. 6 shows SEM images of the amorphous coating in secondary electrons (SE) and back scattered electrons (BSE) modes after 60 min immersion in NaCl solution, just before starting the potentiodynamic polarisation readings. The micrographs confirm that the microcracks are plugged by corrosion products which are seen darker than the matrix in BSE mode (Fig. 6b), because of their lighter weight. Although, the tungsten content and surface morphology of a coating affect the corrosion resistance, but, the amorphisation of coating which happens at higher tungsten content is more likely responsible for the higher corrosion resistance. Generally, the amorphous materials are more corrosion resistant than the crystalline ones on account of the absence of crystallographical defects such as grain boundaries. In spite of the presence of microcracks in amorphous CoeW coating (Fig. 4a), it has still shown a lower corrosion rate compared to the nanocrystalline coating. As was discussed, this is probably because of the microcracks plugging by corrosion products (Fig. 6). Certainly, the microcracks plugging could prevent the build-up of galvanic corrosion cells, and thus eliminate the negative effect of the microcracks on corrosion resistance of the amorphous coating (prevents crevice corrosion).
Table 2 Data extracted from potentiodynamic polarisation plots in Fig. 5. (A: amorphous, N: nanocrystalline, S: substrate). Parameter
Value No motions
E(corr) (mV) i(corr) (mA cm2) ba (V dec1) jbcj (V dec1) OCP (mV) R(p) (U cm2)
Sliding
Ball motion with no touching the specimen
(A)
(N)
(S)
(A)
(N)
(A)
(N)
680 6 0.13 0.06 530 2975
690 8 0.17 0.05 520 2100
170 12 0.07 0.27 150 2015
710 14 0.08 0.08 533 1240
720 18 0.09 0.09 522 1090
700 12 0.09 0.07 532 1425
710 15 0.09 0.09 521 1305
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N. Fathollahzade, K. Raeissi / Materials Chemistry and Physics xxx (2014) 1e10
5
Fig. 6. SEM images of amorphous CoeW electrodeposited coating after 60 min immersion in 3.5 wt-% NaCl solution (before potentiodynamic polarisation test) obtained by detection of (a) secondary (SE) and (b) back-scatter (BSE) electrons.
3.3. Tribocorrosion experiments at OCP
(a)
-520
Potential (mV/Ag/AgCl)
OCP values achieved during the tribocorrosion experiments are related not only to the electrochemical condition of working electrode [22], but also depend on ratio of worn to unworn areas and mass transport condition built up during the sliding [28]. In passive metals, cyclic repassivation/depassivation is also an agent which determines OCP in tribocorrosion experiments. Fig. 7 shows OCP values of the amorphous and nanocrystalline coatings before, during and after ball motion (with and without contact with the working electrode surface). As is seen in Fig. 7, the ball motion (in both contact conditions) causes a significant negative drop in OCP for both coatings. The ball motion alone would enhance the dissolution of the passive film covering the coatings surface [28], if its dissolution were under mass-transport control [29]. It is well known that the thickness of passive films is determined by the
-530
Unloading the ball
-540
(2)
Start sliding
-550
(1)
-560
End sliding
-570 -580 0
Potential (mV/Ag/AgCl)
(b)
1000
2000 3000 Time (s)
4000
5000
-510 Unloading the ball
-520
(2)
-530 Start sliding
-540 -550
End sliding
-560
(1)
-570 0
1000
2000 3000 Time (s)
4000
5000
Fig. 7. Evolution of OCP values of (a) amorphous and (b) nanocrystalline CoeW coatings before, after and during tribocorrosion experiments (1) with and (2) without touching the specimen surface in 3.5 wt-% NaCl solution.
competitive effects of applied potential and film dissolution rate [30]. Thus, by increasing the dissolution rate of passive film due to the mass transport effect (as will be evidenced in Section 3.6), the film becomes thinner and OCP drops to negative values as is seen in Fig. 7. In the case of sliding, OCP is determined not only by the mass transport, but also by some other effects such as: (1) degree of annihilation of passive layer (depassivation) by sliding and extent of active area in worn track, (2) formation of microcracks in the vicinity of wear scar due to the brittleness of coatings, which may enhance corrosion. During the sliding, OCP value gradually became more negative for the amorphous coating (Fig. 7a), whilst this value slightly increased just by ball motion (with no contact). For the nanocrystalline coating (Fig. 7b), OCP slightly increased during the ball motion in both contact conditions. At the end of sliding, OCP values became nobler but they did not approach their initial values. Finally, a significant decrease in OCPs could be observed after unloading the ball from the specimen surface. This decrease is more likely due to the exposure of a very small surface beneath the ball to the solution, which acts more active than other areas and exaggerates the galvanic corrosion. According to Fig. 7a, OCP value of the amorphous coating gradually becomes more active when sliding is going on. This is more probably due to the higher trend of increase in wear scar width of the amorphous coating during the sliding, which surly extends the active area. This more likely arises due to the lower microhardness of the amorphous coating compared with the nanocrystalline one. For the nanocrystalline coating, the variation of OCP during the sliding is very similar to that observed just by ball motion (Fig. 7b) denoting that the total active area in wear scar is likely remained near constant during the sliding. This more possibly arises owing to the higher microhardness of the nanocrystalline coatings. At the end of sliding, restoration of the passive film took place on the worn areas, and hence OCP values gradually became nobler. 3.4. Tribocorrosion experiments under anodic and cathodic polarisation Fig. 8 shows current density readings under anodic polarisation. At the start of applying the anodic potential, the current density of both amorphous and nanocrystalline coatings decreased markedly owing to the thickening of passive layer. It is well known that the electrochemical conditions required for forming passive films are more pronounced at the anodic polarisation than OCP condition. Moreover, under anodic polarisation, the electron extraction force is so great and hence electrons are extracted not only from the conduction band but also from the valence band [31]. In this case, an inversion layer is formed in the space-charge region and the
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N. Fathollahzade, K. Raeissi / Materials Chemistry and Physics xxx (2014) 1e10
(a)
400 (1)
Unloading the ball
200 Start sliding
100
(2)
0 0
1000
(b) 800 Current density (μA cm-2)
Current density (μA cm-2)
300
End sliding
2000 3000 Time (s)
4000
600 500
(1)
Start sliding
400
Unloading the ball
(2)
300 200 0
1000
2000 3000 Time (s)
4000
5000
Start sliding
-20 Unloading the ball
(2)
-40 -60
End sliding
(1)
-80 0
5000
End sliding
700
0
(b)
0
Current density (μA cm-2)
Current density (μA cm-2)
(a)
-20
1000
2000 3000 Time (s)
Start sliding
-40
4000
5000
Unloading the ball
(2)
-60 -80
(1)
End sliding
-100 0
1000
2000 3000 Time (s)
4000
5000
Fig. 8. Evolution of current density readings of (a) amorphous and (b) nanocrystalline CoeW coatings under anodic polarisation (E ¼ þ0.1 V vs. OCP) before, after and during tribocorrosion experiments (1) with and (2) without touching the specimen surface in 3.5 wt-% NaCl solution.
Fig. 9. Evolution of current density readings of (a) amorphous and (b) nanocrystalline CoeW coatings under cathodic polarisation (E ¼ 0.7 V/Ag/AgCl) before, after and during tribocorrosion experiments (1) with and (2) without touching the specimen surface in 3.5 wt-% NaCl solution.
oxidation process occurs more quickly which encourages the growth of passive film [31]. Subsequently, thickening of the passive film causes a decrease in the conductivity of oxide layer, an increase in the band gap, and thus, a significant decrease in the rate of oxidation process. This stabilizes the current density and it remains constant until the ball motion is started. By the ball motion, a sharp increase in current density was observed due to mass transport effect in non-contact condition and also depassivation in sliding with contact condition. However, the current density decreased rapidly to its initial values when the ball motion finished. By unloading the ball from the specimen surface, a slight increase in current density followed by a gradual decrease was observed for both amorphous (Fig. 8a) and nanocrystalline (Fig. 8b) coatings. Under anodic polarisation, the decrease in passive film thickness may be responsible for appearance of a peak at the start of ball motion in non-contact condition (Fig. 8a and b). Over the course of sliding, owing to the gradual decrease of OCP for the amorphous coating (Fig. 7a), it should experience higher anodic polarisation, and thus, thickening and restoration of passive film is expected. The minor alteration in anodic current of the amorphous coating under anodic polarisation (Fig. 8a) indicates that the passive film formed on the surface of the amorphous coating is more protective than that formed on the nanocrystalline one. Thus, the effect of sliding (or ball motion alone) on electrochemical degradation is less intense for the amorphous coating, because lower anodic currents were achieved (Fig. 8a and b). Fig. 9 shows current density readings under cathodic polarisation. Under cathodic polarisation, the passive film can be totally (or partially) reduced from the coatings surface. Moreover, no significant corrosion attack is anticipated because of the cathodic protection of specimen. This fact could be observed in Fig. 9, where
the current density drastically approaches very low values by applying the cathodic potential. By the sliding, the current density became more cathodic probably due to the acceleration of cathodic reactions such as hydrogen evolution on the wear scars and also mass transport effect. At the end of sliding, the current density returned to its initial magnitude in both contact conditions. Again, by unloading the ball from specimen surface, an increase in current density followed by a gradual decrease is observed for both amorphous (Fig. 9a) and nanocrystalline (Fig. 9b) coatings. In this case, the ball motion eliminates the concentration polarisation, and hence the rate of the hydrogen evolution reaction (or the reduction of water) goes up. Because of the absence of considerable corrosion reaction at cathodic polarisation, volume loss of the coatings is mainly mechanical wear and proportion of corrosion accelerated wear (CAW) is negligible. Indeed, there is an inverse relation between the wear rate and hardness of the coatings at cathodic condition [30]. During the sliding, the recorded anodic current density, ia, in Fig. 8 could be achieved by: ia ¼ iw þ ip
(3)
where iw is the current density arises from the ball motion and ip is the anodic current density in the absence of sliding. Furthermore, iw can be considered as: iw ¼ iMT þ iWAC
(4)
where iMT and iWAC are the current densities originated from mass transport effect (MT) and wear accelerated corrosion (WAC), respectively. Table 3 summarizes the proportion of WAC to MT for
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N. Fathollahzade, K. Raeissi / Materials Chemistry and Physics xxx (2014) 1e10 Table 3 The ratio of wear accelerated corrosion (WAC) to mass transport (MT) phenomenon in evolution of tribocorrosion data plotted in Figs. 7e9. Coating structure
Conditions OCP
Amorphous Nanocrystalline
Z
Cathodic polarisation
WAC
MT
WAC
MT
WAC
MT
9% 4%
91% 96%
22% 17%
78% 83%
18% 14%
82% 86%
Z ia dt ¼
Z iMT dt þ
qa ¼ qMT þ qWAC þ qp
provides stronger hydrogen evolution probably due to the induced stresses of cold-working [32]. 3.5. Characterisation of wear scars at OCP conditions
Anodic polarisation
the coatings in sliding at OCP, anodic and cathodic conditions. These values were calculated using the integration of current density vs. time at sliding time intervals. For example in the case of anodic polarisation (curve (1) in Fig. 8), the total charge, qa, is given by:
qa ¼
7
Z iWAC dt þ
ip dt
(5)
(6)
In other words, qa is calculated by integration of curve (1) in Fig. 8. The total charge in no contact condition, qb, is the sum of mass transport charge, qMT, and the anodic charge in sliding condition, qp, and is calculated by integration of curve (2) in Fig. 8. However, the anodic charge, qp, is found from Faraday's low. On the other hand, the charge of WAC, qWAC, equals qa qb. At OCP condition, the ratios in Table 3 were determined by averaging the OCP values during the sliding. The higher proportion of mass transport in tribocorrosion experiments (Table 3) implies that the dissolution reaction of the coatings is mainly a mass-transport controlled process. In all conditions, the proportion of WAC of the amorphous coating is greater than that of the nanocrystalline one. A possible reason is lower hardness of the amorphous coating compared to the nanocrystalline one, which is responsible for its higher damage. Under anodic polarisation, a higher damage is revealed by a wider wear scar. On the other hand, under cathodic polarisation, a higher damage
Line profiles and SEM observations of wear scars after the tribocorrosion experiments give useful data in order to determine the shape and width of the wear scars for calculating total volume loss of the working electrode. Figs. 10 and 11 respectively show SEM images and line profiles of the wear scars created at OCP condition for the amorphous and nanocrystalline CoeW coatings. The width and depth of the wear scars in the amorphous coating were estimated to be about 390 and 4.5 mm, respectively, whilst they were 340 and 3.3 mm for the nanocrystalline coatings. Considering the shape of wear scar as a part of a cylinder, the total volume loss of the amorphous and nanocrystalline coatings was estimated to be around 6 103 and 4 103 mm3, respectively. The higher total volume loss of the amorphous coating is in agreement with its lower microhardness compared to the nanocrystalline coating. The lower microhardness of the amorphous coating results higher percentage of WAC compared to the nanocrystalline one. After the tribocorrosion experiment, some microcracks were detected in the vicinity of the wear scars even in the nanocrystalline coating (Fig. 10b), where no microcracks were observed before sliding (Fig. 4b). The formation of these superficial microcracks in the vicinity of wear scars might be another reason for the difference in OCP values observed at the end of tribocorrosion experiments comparing its initial values (Fig. 7). These microcracks were extended or created by applying the load due to the coatings brittleness. According to SEM images of the coatings (Fig. 10), it seems that the generated microcracks are superficial. Thus, for both coatings, a surface fatigue wear is associated (Fig. 10). According to Fig. 10, the aggregation of wear debris in environs of the wear scars in the amorphous coating is more than that of the nanocrystalline one. 3.6. EIS measurements Fig. 12 shows Nyquist and Bode plots of the amorphous and nanocrystalline CoeW coatings before and after the tribocorrosion
Fig. 10. SEM images (BSE) of wear scars of (a) amorphous and (b) nanocrystalline coatings after the tribocorrosion experiments at OCP condition.
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N. Fathollahzade, K. Raeissi / Materials Chemistry and Physics xxx (2014) 1e10
Fig. 11. Line profiles of wear scars of amorphous (A) and nanocrystalline (N) CoeW coatings (Reconstructed Figure).
experiments at OCP. The appearance of two time constants could be easily distinguished from the plots. If the coatings have pores and cracks, they can also interfere in corrosion process lead to a capacitive response at low frequencies [33]. In Nyquist plots (Fig. 12), the high-frequency loop (the smaller one) is the response of passive film/electrolyte interface, whilst the low-frequency loop (the larger one) is connected to the coating defects and arises due to a finite-length diffusion process of reactants inside the coating defects [34]. The electrical elements extracted from Nyquist plots using the equivalent circuit model depicted in Fig. 13 are listed in Table 4. In the equivalent circuit, R(s) denotes the uncompensated solution resistance which depends upon the distance between reference and working electrodes and solution conductivity. The potentials of points A, Y and B (Fig. 13) represent the potentials at the reference electrode, solution side of double layer and coating surface,
respectively. In order to simulate the effect of passive layer at high frequencies, R(ox) and CPE(ox) were applied which correspond to the oxide resistance and the constant phase element (CPE), respectively. In order to represent the deviation of a capacitive loop from the ideal capacitive behaviour (i.e. depression), it is common to replace the capacitors with CPEs. According to the literature [35], this deviation has been related to the heterogeneities of surface or non-uniform current distribution. The higher R(ox) of the passive film formed on the amorphous coating at OCP (Table 4) confirms the presence of a less defective passive film. On the other hand, R(d) and CPE(d) are equivalent to the diffusion resistance and CPE of diffusion process, which are connected to the low-frequency loop [34]. As is seen in Table 4, a smaller value of CPE(d) Power (0.7595) is obtained for the amorphous coating before sliding. Hence, the centre of its lowfrequency loop fell lower than that for the nanocrystalline coating (Fig. 12a). This more probably arises due to the presence of the microcrack network in the amorphous coating affecting the diffusion process. EIS experiments also provide a complete study on the role of intermediate species adsorbed on a corroding surface [36]. Typically, the magnitude of CPE(ox) should be in order of microFarads per square centimetre [37], but as is seen in Table 4, it has increased significantly to a higher order, particularly before tribocorrosion tests. The main reason is most likely the participation of the adsorbed intermediates in the interfacial corrosion process [33] (e.g. Co(OH)ads in cobalt dissolution process). However, the order of CPE(d) is milliFarads per square centimetre as supposed for finite-length diffusion processes [37]. The polarization resistance, R(p), is the sum of resistances and defined as the impedance limit at zero frequency found by replacing the capacitors (or CPEs) with open circuits [38]. Table 4 reveals that before and after tribocorrosion experiments, R(p) values of the amorphous coating are higher than those of the nanocrystalline coating. As R(p) is inversely related to the corrosion rate (Eq. (2)), the corrosion resistance of the amorphous coating is higher than that of the nanocrystalline one. This is in accordance to
(a) 500
90 Fit result (A)
Phase angle (deg)
-Zimg (Ω cm2)
400
(N)
Fit result (N)
300
(A)
200 100 0
0
(A)
50 30
(N)
10 -10 -30
20 40 60 80 100 120 140 160 180
0.1
Zre (Ω cm2)
(b) 700
10
1000
100000
Frequency (Hz) 70
600
Fit result (A)
500
Fit result (N)
Phase angle (deg)
-Zimg (Ω cm2)
70
400 300 200 100 0 0
150
300
450
600
Zre (Ω cm2)
750
900
50 30 10 -10 -30 0.1
10
1000
100000
Frequency (Hz)
Fig. 12. Nyquist (on the left) and Bode (on the right) plots of the amorphous (A) and nanocrystalline (N) CoeW coatings (a) before and (b) after tribocorrosion experiments at OCP condition in 3.5 wt-% NaCl solution.
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N. Fathollahzade, K. Raeissi / Materials Chemistry and Physics xxx (2014) 1e10
Fig. 13. The equivalent circuit applied for EIS plots in Fig. 12. (R: resistor, CPE: constant phase element).
the results obtained by potentiodynamic polarisation readings and the values of R(p) obtained from EIS plots are very similar to those found from polarization curves (Table 2). According to Table 4, R(ox) of both coatings has decreased after sliding. As R(ox) stems from the entire surface of test specimen (worn and unworn areas), the decrease of R(ox) means that the passive layer formed in worn area after the end of sliding, has not yet regained its initial thickness and/or protective properties [39]. 4. Conclusions In this investigation, the amorphous and nanocrystalline CoeW coatings with about 15 mm thickness were electrodeposited on copper substrate. The tungsten content of the nanocrystalline coating increased from ~35 to ~48 wt-% by increasing the deposition current density and bath temperature where the amorphous coating was obtained. The grain size of the nanocrystalline CoeW coating was estimated to be about 15 nm. A nodular morphology was exhibited for both amorphous and nanocrystalline CoeW coatings, but a microcrack network was only observed in the amorphous CoeW deposit. Briefly, it could be concluded that: 1) The amorphous coating was more resistant to corrosion attacks in 3.5 wt-% NaCl solution compared to the nanocrystalline one and the microcracks plugging could prevent the build-up of galvanic corrosion cells, and thus eliminate the negative effect of the microcracks on corrosion resistance of the amorphous coating. 2) During OCP sliding, the potential of amorphous coating gradually became more active probably due to the widening of wear scar, and thus expansion of active area. This is attributed to the lower microhardness of amorphous coating compared with the nanocrystalline one. 3) Tribocorrosion behaviour of CoeW coatings at OCP could be explained in terms of mass transport (MT) and wear accelerated corrosion (WAC). The proportion of MT in tribocorrosion was higher than WAC, which implied that the corrosion reaction of the coatings was mainly a mass transport controlled process.
Table 4 Data extracted from EIS plots in Fig. 12 using the equivalent circuit in Fig. 13. Element
Value Before tribocorrosion
After tribocorrosion
Amorphous Nanocrystalline Amorphous Nanocrystalline R(s) (U cm2) CPE(ox) T (F cm2) CPE(ox) Power R(ox) (U cm2) CPE(d) T (F cm2) CPE(d) Power R(d) (U cm2) R(p) (U cm2)
1.57 0.002388
1.88 0.005291
4.83 0.001206
12.74 0.001254
0.6435 40.18 0.02289
0.5479 19.43 0.02477
0.3622 7.28 0.001155
1.0000 8.39 0.002893
0.7595 2940 2980
0.9000 2090 2110
0.8661 2447 2454
0.6095 1835 1843
9
3) In anodic sliding, the evolution of current readings indicates that the effect of sliding (or ball motion alone) on electrochemical degradation is more intense for the nanocrystalline coating, because of the higher achieved anodic currents. In this case, the passive film formed on the surface of the amorphous coating was more protective than that formed on the nanocrystalline one. However, the volume loss of the amorphous coating was higher than that of the nanocrystalline one. This is related to the higher microhardness of nanocrystalline coating. 4) Under cathodic polarisation, the passive film could be totally (or partially) reduced in tribocorrosion experiment, but no remarkable corrosion was observed. In this case, the ball motion eliminated the concentration polarisation, and hence the rate of the hydrogen evolution reaction (or the reduction of water) went up. References [1] Z. Ghaferi, K. Raeissi, M.A. Golozar, H. Edris, Characterization of nanocrystalline CoeW coatings on Cu substrate, electrodeposited from a citrate-ammonia bath, Surf. Coat. Technol. 206 (2011) 497e505. [2] P. Lammel, L.D. Rafailovic, M. Kolb, K. Pohl, A.H. Whitehead, G. Grundmeier, B. Gollas, Analysis of rain erosion resistance of electroplated nickeletungsten alloy coatings, Surf. Coat. Technol. 206 (2012) 2545e2551. [3] L. Wang, Y. Gao, T. Xu, Q. Xue, A comparative study on the tribological behaviour of nanocrystalline nickel and cobalt coatings correlated with grain size and phase structure, Mater. Chem. Phys. 99 (2006) 96e103. [4] F. Su, C. Liu, P. Huang, Establishing relationships between electrodeposition techniques, microstructure and properties of nanocrystalline CoeW alloy coatings, J. Alloys. Compd. 557 (2013) 228e238. [5] F. Su, P. Huang, Microstructure and tribological property of nanocrystalline CoeW alloy coating produced by dual-pulse electrodeposition, Mater. Chem. Phys. 134 (2012) 350e359. [6] D.P. Weston, S.J. Harris, H. Capel, N. Ahmed, P.H. Shipway, J.M. Yellup, Nanostructured CoeW coatings produced by electrodeposition to replace hard Cr on aerospace components, Trans. Inst. Met. Finish. 88 (2010) 47e56. [7] H. Capel, P.H. Shipway, S.J. Harris, Sliding wear behaviour of electrodeposited cobaltetungsten and cobaltetungsteneiron alloys, Wear 255 (2003) 917e923. [8] D.P. Weston, P.H. Shipway, S.J. Harris, M.K. Cheng, Friction and sliding wear behaviour of electrodeposited cobalt and cobaltetungsten alloy coatings for replacement of electrodeposited chromium, Wear 267 (2009) 934e943. [9] F. Su, C. Liu, P. Huang, Friction and wear of nanocrystalline Co and CoeW alloy coatings produced by pulse reverse electrodeposition, Wear 300 (2013) 114e125. [10] A. Igual Munoz, N. Espallargas, Tribocorrosion mechanisms in sliding contacts, in: D. Landolt, S. Mischler (Eds.), Tribocorrosion of Passive Metals and Coatings, Woodhead Publishing Limited, 2011, p. 118. [11] P. Jemmely, S. Mischler, D. Landolt, Electrochemical modeling of passivation phenomena in tribocorrosion, Wear 237 (2000) 63e76. [12] E. Arslan, Y. Totik, I. Efeoglu, The investigation of the tribocorrosion properties of DLC coatings deposited on Ti6Al4V alloys by CFUBMS, Prog. Org. Coat. 74 (2012) 768e771. [13] N. Sato, T. Ohtsuka, Anodic oxidation of cobalt in neutral and basic solution, J. Electrochem. Soc. 125 (1978) 1735e1740. [14] M. Metikos-Hukovic, R. Babic, Passivation and corrosion behaviours of cobalt and cobaltechromiumemolybdenum alloy, Corros. Sci. 49 (2007) 3570e3579. [15] M. Pontinha, S. Faty, M.G. Walls, M.S. Ferreira, M. Da Cunha Belo, Electronic structure of anodic oxide films formed on cobalt by cyclic voltammetry, Corros. Sci. 48 (2006) 2971e2986. [16] H. Jung, A. Alfantazi, Corrosion properties of electrodeposited cobalt in sulfate solutions containing chloride ions, Electrochim. Acta 55 (2010) 865e869. [17] A.M. Human, B. Roebuck, H.E. Exner, Electrochemical polarisation and corrosion behaviour of cobalt and Co(W,C) alloys in 1 N sulphuric acid, Mater. Sci. Eng. A 241 (1998) 202e210. [18] A. Bodaghi, J. Hosseini, Corrosion behavior of electrodeposited cobaltetungsten alloy coatings in NaCl aqueous solution, Int. J. Electrochem. Sci. 7 (2012) 2584e2595. [19] Z. Abdel Hamid, Electrodeposition of cobaltetungsten alloys from acidic bath containing cationic surfactants, Mater. Lett. 57 (2003) 2558e2564. [20] R.C. Santana, A.N. Campos, E.A. Medeiros, A.M. Oliveira, L.F. Oliveira, S. Prasad, Studies on electrodeposition and corrosion behaviour of a NieWeCo amorphous alloy, J. Mater. Sci. 42 (2007) 9137e9144. [21] H.J. Yun, S. Dulal, C.B. Shin, C. Kim, Characterisation of electrodeposited CoeWeP amorphous coatings on carbon steel, Electrochim. Acta 54 (2008) 370e375. [22] S.K. Ghosh, J.P. Celis, Tribological and tribocorrosion behaviour of electrodeposited CoW alloys and CoWeWC nanocomposites, Tribol. Int. 68 (2013) 11e16.
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Please cite this article in press as: N. Fathollahzade, K. Raeissi, Electrochemical evaluation of corrosion and tribocorrosion behaviour of amorphous and nanocrystalline cobaltetungsten electrodeposited coatings, Materials Chemistry and Physics (2014), http://dx.doi.org/ 10.1016/j.matchemphys.2014.07.013