Effect of electrochemical corrosion on the subsurface microstructure evolution of a CoCrMo alloy in albumin containing environment

Applied Surface Science 406 (2017) 319–329

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Effect of electrochemical corrosion on the subsurface microstructure evolution of a CoCrMo alloy in albumin containing environment Zhongwei Wang, Yu Yan ∗ , Yanjing Su, Lijie Qiao Corrosion and Protection Center, Key Laboratory for Environmental Fracture (MOE), University of Science and Technology Beijing, 100083, China

a r t i c l e

i n f o

Article history: Received 14 December 2016 Received in revised form 16 February 2017 Accepted 17 February 2017 Available online 20 February 2017 Keywords: Cobalt Bio-tribocorrosion AES EIS TEM Interfaces

a b s t r a c t The subsurface microstructures of metallic implants play a key role in bio-tribocorrosion. Due to wear or change of local environment, the implant surface can have inhomogeneous electrochemical corrosion properties. In this work, the effect of electrochemical corrosion conditions on the subsurface microstructure evolution of CoCrMo alloys for artificial joints was investigated. Transmission electron microscope (TEM) was employed to observe the subsurface microstructures of worn areas at different applied potentials in a simulated physiological solution. The results showed that applied potentials could affect the severity of the subsurface deformation not only by changing the surface passivation but also affecting the adsorption of protein on the alloy surface. © 2017 Elsevier B.V. All rights reserved.

1. Introduction Due to the excellent mechanical, corrosion-resistant, wearresistant and biocompatible properties, CoCrMo alloys are widely used as orthopaedic implants for artificial joints [1–3]. Nevertheless, they would suffer from bio-tribocorrosion via joint motions that caused by the combined effect of wear and corrosion in biological environments [4]. The passive film formed on a CoCrMo alloy surface can be damaged or even removed in the wear process, resulting in accelerated metal dissolution which in turn can lead to accelerated wear in many cases [5]. The generation of wear debris and metallic ions induced by bio-tribocorrosion is the biggest issue, and could cause periprosthetic tissue reaction, osteolysis, late aseptic loosening and long-term toxicological effects [6–8]. This issue have been a serious threat to the safe and long-term use of artificial joints. The subsurface microstructure of metallic materials plays a key role in their wear and electrochemical corrosion properties. It can affect debris properties and the amount, the wear rate, the corrosion rate and the character for passivation etc. [9–11]. Under tribological contact, the subsurface microstructure can change during the plastic deformation process, such as the appearance of

∗ Corresponding author at: Corrosion and Protection Center, University of Science and Technology Beijing, No. 30 Xue Yuan Road, Room 413, Beijing 100083, China. E-mail addresses: [email protected], [email protected] (Y. Yan). http://dx.doi.org/10.1016/j.apsusc.2017.02.152 0169-4332/© 2017 Elsevier B.V. All rights reserved.

defects, the decrease in grain size and even nanocrystallization [12,13]. These changes can in turn affect the tribological behaviour of the materials. For example, the introduction of defects, such as dislocations, twins and stacking faults, can enhance the hardness and strength of the material surface and decrease its wear rate [14,15]. A decrease in grain size can also increase the hardness and strength of the material’s surface by blocking the movement of any dislocations with the high grain boundary density [16]. This can enhance the wear-resistance of the surface but decrease the size of debris size and increase the amount of debris that may accelerate the failure of artificial joints [11]. The subsurface microstructure could also affect the subsequent subsurface deformation. Zeng et al. reported that the nanocrystalline layer formed on the CoCrMo alloy topmost surface could induce the deformation beneath it at a low strain stress [17]. Thus, it could be known that the subsurface microstructures of CoCrMo alloys exhibit a close relationship to the long-term and safe service of artificial joints. Electrochemical corrosion is one of the main aspects of biotribocorrosion. It can determine the surface’s chemical composition and passivation condition. The metallic materials would show certain values of electrode potential after being implanted in the body. However, the implant material cannot maintain a uniform electrode potential across the whole surface. For example, in the artificial joint, the rubbed zone which has lost the protection of the passive film can acquire a lower potential when engaged in galvanic coupling with the surrounding area. Furthermore, the surrounding area can yield a higher potential [18]. The local environment,

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Fig. 1. Dynamic potential polarization curve of the CoCrMo alloy in the phosphate buffer solution with bovine serum albumin during tribological contact.

such as the pH, proteins and coupling materials, can also change the electrode potential of the implanted materials [19,20]. These changes would affect the metal surface properties. Fox example, the applied potential can affect the adsorption of proteins on medical implant metal [21]. Furthermore, the electrochemical corrosion can also affect the plastic deformation of the subsurface. More deformation in stainless steel was found with passively applied potential, and this effect was tentatively explained by the interaction of the dislocation and the passive film in previous studies [22–24]. The presence of the passive film also significantly increased subsurface deformation in a Stellite CoCrMo alloy in the sulphuric acid aqueous solution [25]. As might be expected, the electrochemical corrosion would change the subsurface evolution of biomedical CoCrMo alloy during bio-tribocorrosion in the body not only by affecting the passivation but also its interaction with physiological environment. However, there are few studies concerned with this area. The aim of this study is to investigate the subsurface microstructures of CoCrMo alloy after worn at different electrochemical corrosion conditions in vitro. This can be helpful in understanding the failure mechanism and can guide the design of implanted materials and artificial joints. In this study, the CoCrMo alloys were worn within a simulated physiological solution with different applied potentials standing for different electrochemical corrosion conditions. The subsurface microstructures were observed on the cross-section samples from the worn surfaces. The adsorption of proteins on CoCrMo alloy surfaces with different applied potentials was also investigated.

water and ultrasonically cleaned in ethanol and finally dried in the air. The experimental tests were carried out in phosphate buffered saline (PBS) solution with 10 g L−1 bovine serum albumin (BSA). The composition of the PBS contained 8 g L−1 NaCl, 0.2 g L−1 KCl, 1.44 g L−1 Na2 HPO4 and 0.25 g L−1 KH2 PO4 . The pH value was 7.4. All the tests were carried out at 37 ± 0.5 ◦ C in the atmospheric environment. 2.2. Electrochemical measurements

2. Experiments

The dynamic potential polarization was used to select the electrochemical properties of the CoCrMo alloy using an electrochemical work station (CH660E, CH Instruments). The scanning rate was 1 mV s−1 . Fig. 1 shows the dynamic potential polarization curve of the CoCrMo alloy in the PBS with BSA during tribological contact (the parameters are same with that in tribological tests and showed in the next section). Three typical electrochemical corrosion conditions were selected: −0.8 VAg/AgCl for the cathodic protection condition, the open circuit potential (OCP) for the normal condition, and 0.2 VAg/AgCl for the passive condition. EIS was used to investigate the effect of the applied potential on the BSA adsorption of the bulk CoCrMo alloy in a static state. EIS measurements were initiated after the as-polished samples were immersed in the PBS with BSA with different applied potentials for 1 h. The EIS measurement ranged from 105 to 10−2 Hz with alternating current amplitude of ±10 mV. The EIS experiments were repeated four times at each potential. The impedance data were analysed with the ZsimpWin V3.10 software.

2.1. Materials

2.3. Tribological tests

Rods of a high-carbon CoCrMo forged alloy specified in ASTM standard F1537-11 were used as the substrate material (Produced by Beijing Rongdian Metallic Materials Co., Ltd., China). The chemical composition was 28.32 wt.% Cr, 5.43 wt.% Mo, ∼0.2 wt.% C and Co balance. The CoCrMo rods were cut into disks of 20 mm in diameter and 5 mm in thickness. The samples were ground with SiC papers of up to 3000 grit and were polished with a diamond paste until they had a mirror-like surface finish (surface roughness Ra around 10 nm). After polishing, the samples were rinsed with deionized

In the tribological tests, the CoCrMo alloy plate was embedded with epoxy resin in a polyvinyl chloride (PVC) tube of 25 mm in diameter and 8 mm in thickness (Fig. 2a), and then connected with a wire and then polished, as mentioned above. In order to acquire a homogeneous and horizontal wear track, the linear contact geometry (cylinder-on-plane) was used in this study. The counter material was the same as the plate. The dimensional drawing is shown in Fig. 2b. The surface pre-treatment of the counterpart was the same as the plate.

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Fig. 2. Schematic diagrams of the tribological testing set up: (a) substrate plate, (b) counterpart and (c) measurement cell.

Fig. 2c shows the diagram of the integrated tribocorrosion test set up. The tribological experiment was using the following steps: Step 1: the sample which was shown in Fig. 2a was fixed in the centre of a teflon cell as the working electrode. A platinum plate and an Ag/AgCl electrode were also fixed on the cell inside as the auxiliary electrode and the reference electrode. Step 2: the three electrodes were connected with an electrochemical workstation. Then the cell was assembled with a tribometer (UMT2, CETR, Bruker). Step 3: 100 mL electrolyte was added into the cell and a selected potential (for 0.2 VAg/AgCl or −0.8 VAg/AgCl ) was applied to the sample under the potentiostatic model or at the OCP condition by using the three electrode set-up. The current values (for 0.2 VAg/AgCl or −0.8 VAg/AgCl ) or potential value (for the OCP) were monitored by the electrochemical workstation. The electrolyte temperature was maintained at 37 ± 0.5 ◦ C throughout the entire test. Step 4: after stabilization for 10 min, the polished counterpart was moved to the plate and a load of 8 N was applied. Based on the Hertzian contact theory, the contact stress on the plate was 231 MPa, initially. Then the cell reciprocated with a frequency of 1.5 Hz while the sliding distance was 30 mm per cycle. Step 5: The rubbing motion was stopped after running for 48 h. The potentiostatic measurement or the OCP measurement was stopped at the same time. Then the sample was taken out and gently rinsed with deionized water and dried in the air for the subsequent tests. The total wear loss of the plate was quantified using a contour graph (Dekatk150, VEECO) by characterizing the surface topography in and around the wear track. Thus the cross sectional area of the wear track would be known. Finally, the total wear loss volume of the plate could be given by the product of cross sectional area of the wear track and length of the wear track. The volume lost of the counterpart was gained by measuring the difference in height before and after the test and then calculating the volume lost. Four pairs of worn plate and pin were tested to gain the statistic values of the volume loss.

2.4. Surface and subsurface characterization A Field Emission Scanning Electron Microscopy (FESEM, SUPRA 55, Zeiss) was used to acquire the surface microscopy image of the wear track from the secondary electron. The microstructure of the cross-section of the as-polished CoCrMo alloy was gained by using a FESEM.

Longitudinal cross-sections of the wear track were prepared for TEM using a focused ion beam (FIB, Helios nanolab 600, FEI). The area of interest was located in the centre of the wear track. A tungsten or platinum deposition layer was deposited on the top surface of the areas of interest before cutting to avoid exposure to the Ga+ beam. After this, the FIB specimens were observed using a TEM (Tecnai F20, FEI) with an accelerating voltage of 200 kV. 2.5. Auger electron spectroscopy The morphologies and chemical compositions of the CoCrMo alloy surface after being immersed in the PBS with BSA at different applied potentials in a static state were acquired using the FESEM and Auger electron spectroscopy (AES, PHI-700, ULVACPHI), respectively. Before the test, the polished sample was fixed in the testing cell and 100 mL PBS with the BSA solution was added into the cell. Then a potential was applied for 6 h without tribological test. The sample was then taken out and gently washed by deionized water to remove the un-adsorbed protein. Finally, the sample was dried in the air and observed with the SEM and AES. The accelerating voltage of the electron-beam was 5 keV in the AES tests. The depth profile acquisition was performed by the Ar+ beam. The sputter rate of the SiO2 standard measured under these conditions (6 nm min−1 ) was used to convert the sputter time to the approximate sputter depth. 3. Results 3.1. Bio-tribocorrosion behaviour Fig. 3a shows the evolution of the coefficient of friction (COF) for the different electrochemical corrosion conditions. The curve at 0.2 VAg/AgCl shows a different evolution when compared with the other two conditions: after a slight decrease, the COF increased rapidly and then increased slightly until the rubbing ceased. The COF at 0.2 VAg/AgCl was higher than for the other two electrochemical conditions in the entire experiment. At OCP and −0.8 VAg/AgCl , the COFs decreased rapidly then increased, and finally slightly decreased until the rubbing ceased. The COF at −0.8 VAg/AgCl was slightly lower than that at OCP. Fig. 3b shows the evolution of the potential with time at the open circuit condition. It can be observed that the OCP shifts drastically towards to the negative direction when the motion started. This could be explained by that the passive film on alloy surface was mechanically damaged or even removed, which could lead to

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Fig. 3. (a) Evolution of the coefficient of friction with time for the different electrochemical corrosion conditions, (b) evolution of the potential with time at the OCP condition and (c) evolutions of currents with time at 0.2 VAg/AgCl and −0.8 VAg/AgCl .

the exposure of a flash metal surface and a decrease of the potential [26]. Then the OCP has a steady decrease and maintains at about −0.45 VAg/AgCl during sliding. When the sliding test was ended, the OCP increases rapidly and this is also induced by the repassivation process of the wear track. Fig. 3c shows the evolution of current values with time at 0.2 VAg/AgCl and −0.8 VAg/AgCl during sliding. A positive current can be observed at 0.2 VAg/AgCl and a negative current value is observed at −0.8 VAg/AgCl . During the entire test at −0.8 VAg/AgCl , no bubble can be observed at the sample surface. Table 1 shows the statistic wear volumes of CoCrMo alloys for different applied potentials. The wear volume decreased with a decrease in the applied potential. In the tribocorrosion process, the total material loss could be divided into four components: the

Fig. 4. Secondary electron images of surface topographies of wear tracks for different electrochemical corrosion conditions.

wear damage in the absence of corrosion, the corrosion loss in the absence of wear, the change in the corrosion rate due to wear and the change in the wear rate due to corrosion. At the cathodic protection potential, the anodic corrosion reactions can be inhibited and the material degradation is only caused by mechanical wear. At OCP condition, the corrosion current was 2.3 × 10−5 A cm−2 (obtained from Fig. 1). With the potential increased to 0.2 VAg/AgCl , the corrosion rate was accelerated to about 5 × 10−3 A cm−2 (Fig. 3c). The electrochemical corrosion activates the later three components and lead to the increase of the volume lost. Fig. 4 shows the surface topographies of wear tracks at different electrochemical corrosion conditions. It can be observed that there were some scratches over the entire scar length as well as parallel to the sliding direction at 0.2 VAg/AgCl . The scratches were not obvious at OCP and −0.8 VAg/AgCl . All the wear-track surfaces were covered

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Table 1 Statistic wear volumes of CoCrMo alloys at different electrochemical corrosion conditions. Four pairs of worn plate and pin were tested to gain the statistic volume loss. Data were presented as the mean ± standard deviation (n = 4). Corrosion condition

Counterpart (×10−3 mm3 )

Flat (×10−3 mm3 )

Total volume lost (×10−3 mm3 )

0.2 VAg/AgCl OCP −0.8 VAg/AgCl

44.9 ± 5.3 25.9 ± 4.9 8.5 ± 1.9

20.4 ± 4.6 18.3 ± 4.0 11.6 ± 2.2

65.3 ± 9.9 44.2 ± 8.9 20.1 ± 4.1

Fig. 5. HAADF TEM images and elements distribution map from EDS of the wear track subsurface at 0.2 VAg/AgCl .

with discontinuous tribo-films (the dark-coloured area). It can be observed that the coverage fraction of tribo-film increased as the applied potential decreased. Fig. 5 shows the element distribution maps from the energydispersive X-ray spectroscopy (EDS) tests for the wear track subsurface at 0.2 VAg/AgCl which includes the substrate, the tribofilm and the protective Tungsten layer from the bottom to the top. Though EDS results are not accurate for the quantitative analysis of carbon due to the low sensitivity and the contaminant issues, the EDS results in this map of carbon were acceptable when treated as qualitative or even semiquantitative result. At this stage, we are interested in the differences between the tribo-film and the substrate. The enrichment of oxygen in the tribo-film may due to the presence of oxides, hydroxides or denatured proteins. However, the metallic elements are almost not detected in the tribo-film. The sharp increase of the carbon content and the appearance of oxygen can indicate that the tribo-film was transformed from the proteins, because the bovine serum albumin has C:O mass ration of 1:0.225 [27]. This kind of the tribo-film was also found on CoCrMo hip joints in vivo and in vitro [27,28].

3.2. Subsurface microstructures Fig. 6 shows the back scattered electron image of the crosssection of the as-polished CoCrMo alloy sample. The microstructure of the CoCrMo alloy consists of an equiaxed structure with a grain size ranging from 10 to 20 ␮m. Spherical carbides embedded randomly in the Co-rich matrix could be observed. The dark oblique lines in the grain were ␧-martensite, which was induced in the forged process by strain-induced phase transformation. There is no obvious difference between the near-surface zone and the deeper zone. This means that no obvious plastic deformation was induced to the subsurface during the polishing process. Fig. 7 shows the TEM bright-filed images of cross-sections of the wear tracks for different electrochemical corrosion conditions.

Fig. 6. Back scattered electron image of the cross section of the as-polished CoCrMo alloy sample.

A severe plastic deformation layer with a thickness about 1.5–2 ␮m can be observed at 0.2 VAg/AgCl from Fig. 7a and d. Many stacking faults were introduced into the subsurface when compared with the deformation-free zone at the bottom of the specimen. The stacking faults of high density can be found easily in the deformed zone near the surface. The surface was covered with a thin tribofilm. Fig. 7b and e shows the subsurface microstructures of the wear tracks for OCP. The subsurface deformation was slighter than that at 0.2 VAg/AgCl . The density of the stacking faults was lower than that at 0.2 VAg/AgCl . The thickness of the severe plastic deformation layer was around 1 ␮m and was also thinner than it. However, the thickness of the tribo-film was much thicker than that at 0.2 VAg/AgCl and was around 400 nm. Fig. 7c and f shows the subsurface microstructures of the wear track at −0.8 VAg/AgCl . The stacking faults (black

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Fig. 7. TEM bright-filed images of cross-sections of the wear tracks for different electrochemical corrosion conditions: (a) (d) 0.2 VAg/AgCl , (b) (e) OCP and (c) (f) −0.8 VAg/AgCl .

area) could also be found from Fig. 7f. But, no obvious severe plastic deformation layer with high density defects was found on the subsurface. The grain size near the surface was much larger than the former two samples. The thickness of the tribo-film increased to around 500 nm at −0.8 VAg/AgCl . Fig. 8 shows the TEM images of the subsurface microstructures near the topmost surface. At 0.2 VAg/AgCl (Fig. 8a and b), there was an ultra-fine grain nanocrystalline (NC) layer. The gain size in this layer was around 5 nm. The NC layer can also be observed at for OCP (Fig. 8c). However, the grain size (50–100 nm) was much larger than that at 0.2 VAg/AgCl . Furthermore, from the selected-area electron diffraction (SAED) image, the NC grains have random orientations. It worth noting that there was a passive film between the tribo-film and the metal substrate. The EDS results show that the passive film mainly consists of O (43.77 at.%), Co (20.37 at.%) and Cr (18.43 at.%). When the applied potential decreased to −0.8 VAg/AgCl , the grain size near the surface was larger than that for 0.2 VAg/AgCl and OCP (Fig. 8d). The SAED result indicates the existence of sub-grains with small misorientation. Under severe strain accumulation, the defect/stacking faults would change to the grain boundary and divide the original grain. Then the sub-grain with small misorientation would form. The summarized subsurface characteristics for different electrochemical conditions are shown in Table 2. The TEM results indicate that the increase in the applied potential can aggravate the subsurface deformation of alloy in bio-tribocorrosion. 3.3. Protein adsorption It was found that proteins could affect the subsurface deformation of the CoCrMo alloy during bio-tribocorrosion in our previous study [29]. Furthermore, Mischler et al. [21] found that protein adsorption on a CoCrMo alloy depended on the applied potential:

the cathodic polarization yielded more BSA adsorption than for passive potentials. However, the experimental conditions in this study exhibit some differences compared with the previous studies, such as the applied potential value and the properties of the substrate, a physical vapour deposition CoCrMo thin-film was used to investigate the BSA adsorption for the cathodic potential in their study. In order to acquire a protein adsorption result which was closer to this study, EIS, SEM and AES were used to investigate the effect of the applied potential on the BSA adsorption on the bulk-CoCrMo alloy in a static state. According to previous works [30–32], it is confirmed that the BSA adsorption on stainless steel and CoCrMo alloy follows the Langmuir isotherm. The surface concentration of the amount of adsorbed protein ( in mol cm−2 ) can be correlated with the charge transfer resistance related to the protein adsorption 1/Rct,c , i.e., 1/Rct,c proportional to  . The 1/Rct,c can be calculated by means of Eq. (1): 1 1 1 = − Rct,c Rct,i Rct,0

(1)

where Rct,i is the charger transfer resistance in PBS with BSA solution and Rct,0 is the transfer resistance recorded in the proteinfree PBS solution. Fig. 9 shows the Bode plots and Nyquist plots of the CoCrMo alloy after the application of different potentials in PBS and PBS with BSA. Two time constants can be observed in the impedance spectra except in the PBS with BSA at −0.8 VAg/AgCl . Characteristic features of Nyquist plot in PBS with BSA at −0.8 VAg/AgCl was the presence of a diffusive contribution at low frequencies, which is called Warburg impedance. This result indicated that the dissolution mechanism of CoCrMo alloy was being controlled by the mass transport rate in the presence of BSA as low potential [33]. The time constant related to the passive film formed

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Fig. 8. TEM and SAED images of the subsurface microstructures near the topmost surface at different electrochemical conditions: (a) (b) 0.2 VAg/AgCl , (c) OCP and (d) −0.8 VAg/AgCl .

Table 2 Summarized subsurface characteristics for different electrochemical corrosion conditions.

Severe deformed layer thickness Stacking faults density in the deformed layer Grain size near the surface Passive film Tribo-film thickness

0.2 VAg/AgCl

OCP

−0.8 VAg/AgCl

1.5–2 ␮m Highest ∼5 nm Unobserved ∼40 nm

∼0.5 ␮m High 50–100 nm Observed ∼400 nm

None Low >100 nm Unobserved ∼500 nm

on the CoCrMo alloy and is attributed to the passive film resistance (Rpass )/capacitance (Cpass ) parallel combination across the passive film. Another one formed by one resistance (Redl )/capacitance (Cedl ) (attributed to the electric double layer) parallel combination. A constant phase angel element (CPE) is introduced to replace capacitor and to account for the non-ideal behaviour of the capacitive elements due to different physical phenomena such as surface heterogeneity. From the fitted EIS data, the charge transfer resistance (Rct ) in different solutions and applied potentials equalled to Rpass and was used for describing the influence of applied potentials on the adsorption of protein on CoCrMo alloy surface. The charge transfer resistance (Rct ) of the process was calculated as the sum of Rpass and Redl. The dependence of the 1/Rct,c (proportional to surface concentration) compared with different electrochemical corrosion conditions is shown in Fig. 10. It can be clearly seen that the 1/Rct,c value increased as the applied potential decreased, especially with changes in the cathodic conditions. The EIS results indicate that the amount of adsorbed protein on the CoCrMo alloy surface will increase as the applied potential decreases. Fig. 11a–c shows the CoCrMo alloy surface morphologies after immersed in PBS with BSA at different electrochemical conditions

for 6 h. It can be observed that the alloy surfaces were covered with various discontinuous films. In Fig. 11c, the EDS results show that the chemical composition of the area covered with the film contains 40.18 wt.% C, 11.31 wt.% O, 30.88 wt.% Co, 14.00 wt.% Cr and 3.36 wt.% Mo. The chemical compositions of the area where was not covered with the film were 10.89 wt.% C, 58.15 wt.% Co, 25.95 wt.% Cr and 5.01 wt.% Mo but O was not detected. This result was expected as, at the cathodic potential, the oxidation reactions were effectively inhibited. The sharp increase of the carbon content and the appearance of O can reveal that the film was transformed from the proteins. Therefore, such films can be formed by an aggregate of adsorbed proteins. On the other side, the coverage fraction of the films increased with the decrease in the applied potential. The surface chemical compositions from different electrochemical corrosion conditions were also investigated using AES and the results are shown in Fig. 11d–f. The carbon element curves indicate that with a decrease of the surface potential, the thickness of the adsorbed protein film can increase. These results agree with the EIS results that the alloy surface could adsorb more protein at lower potential.

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Fig. 9. Bode plots, Nyquist plots and equivalent circuits of CoCrMo alloy after the application of different potentials in phosphate buffer solution and phosphate buffer solution with bovine serum albumin.

The first step is the reduction of proton or water and the formation of adsorbed hydrogen atom on metal surface [34]: H+ + e− → Hads

(2)

or 2H2 O + 2e− → 2Hads + 2OH−

(3)

Then two adsorbed hydrogen atom could combine to form gaseous hydrogen: 2Hads → H2 ↑ (gas)

(4)

In this study, the adsorbed hydrogen atom would participate in the reduction reaction of the BSA [35]: RSSR + 2Hads → RSHHSR ads Fig. 10. The 1/Rct,c values of CoCrMo alloy after the application of different potentials. Data were presented as the mean ± standard deviation (n = 4).

4. Discussion 4.1. Electrochemical reactions The possible reactions that may occur at −0.8 VAg/AgCl in the present solution are the oxygen, water, protons and the BSA reduction [21]. Althought this potential is below the hydrogen equilibrium potential at pH 7.4 (about −0.66 VAg/AgCl ), no bubble could be noticed on the alloy surface during the entire test. This phenomenon could be explained by that the hydrogen atoms from the reduction reaction of proton and water participate in the reduction reaction of BSA. It is well understood that the formation of gaseous hydrogen on the metal surface at the catholic condition is a two-step process.

(5)

where SS—stands for the disulphide bonds in BSA. Thus the adsorbed hydrogen atoms are consumed in Reaction (5) that inhibits the formation of gaseous hydrogen. Moreover, the low concentration of protons at pH 7.4 for Reaction (2), relatively the small over-potential for Reaction (3) and abundant BSA in Reaction (5) would make this inhibition effect more apparent. In particular, the additional disulphide bonds in the BSA will decrease after longer polarization below −0.6 VAg/AgCl , leading to the formation of insoluble products by the following reaction [35]: RSSR + 2H+ + 2e− → RSHHSR

(6)

This might be the reason that the CoCrMo alloy surface can adsorb most BSA (Fig. 11c) and form the thickest tribo-film with the highest coverage (Fig. 4a) at −0.8 VAg/AgCl . The OCP value was around −0.45 VAg/AgCl for sliding. In this case, the cathodic reactions will be inhibited. The cathodic reactions will be severely inhibited and the oxidation reaction will be accelerated at 0.2 VAg/AgCl [36]. The inhibition of BSA reduction and the acceleration of passivation (metallic elements’ oxidation) will decrease the

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Fig. 11. (a–c) Secondary electron images of the CoCrMo alloy surface morphologies and (d–f) surface chemical compositions after the application of different potentials for 6 h in the phosphate buffer solution with bovine serum albumin.

Fig. 12. Diagram for the interaction between the dislocations’ movement and the passive film.

amount of adsorbed protein and inhibit the formation of a tribo-film on the CoCrMo alloy at high potential. 4.2. Effect of electrochemical corrosion conditions on subsurface microstructures From the above experimental results, it can be seen that the applied potential has an obvious effect on the subsurface microstructures’ evolution of a CoCrMo alloy in the biotribocorrosion processes. The active mechanism of the applied potential on the microstructures’ evolution can be explained from the following two sides: 4.2.1. Effect of surface passivation Dislocations generated at the surface can either diffuse towards the bulk of the metal or annihilate when emerging on the metal’s surface during the wear process [24]. Due to the low stacking fault energy, the stacking fault activities dominated the deformation process of the CoCrMo alloy. The dominant mechanisms governing plastic deformation and microstructural refinement include the formation of stacking faults, grain subdivision and martensite transformation as well as the formation of NC [14]. Fig. 12 shows the diagram for the interaction between the dislocations’ movement and the passive film. The presence of a passive film for a high applied potential can reduce or even suppress the dislocations’

annihilation by blocking the metal surface. Thus, the flux of the dislocations towards the bulk material can be enhanced, leading to a higher stacking fault density in the subsurface [23]. Furthermore, the inhibition of annihilation of the emerging dislocations can lead to the accumulation of strains on the subsurface. Such strain accumulation can benefit the formation sub-grains and nanocrystals. Thus, the severest deformation is found at 0.2 VAg/AgCl . It is worth noting that the passive film was only found at OCP. Intuitively, the passive film will be thicker for a passive potential of 0.2 VAg/AgCl than for OCP. However, the passive film is barely found for 0.2 VAg/AgCl (Fig. 8b). This phenomenon might be explained by the protective effect of the tribo-film. The tribo-film can protect the passive film from mechanical scaping. Thus, the passive film can be thick and easily found for OCP. Although the passive rate for 0.2 VAg/AgCl is higher than that for OCP, the amount of adsorbed protein is less than later on due to the inhibition of the cathodic reactions. With the decrease of the number of adsorbed proteins, the thickness and coverage fraction of the tribo-film decrease. This can weaken the protective effect of the tribo-film. Thus, the passive film can easily spall off surfaces and is too thin to observe. The anodic reactions are inhibited at the cathodic potential, thus the passive film is also not found for −0.8 VAg/AgCl . 4.2.2. Effect of the protein adsorption Unlike in normal corrosive environments, such as simple acid or saline solutions, the applied potentials would not only affect the surface passivation but also interact with bio-molecules in the biological environment. In our previous study [29], it was proved that the protein could affect the subsurface microstructure evolution during the tribocorrosion process by transforming to the tribo-film covered the metal surface. In this study, the applied potential was found to affect protein adsorption and change the thickness and coverage fraction of the tribo-film. This effect can be found in the evolution of COF in Fig. 3a. The decrease in the COF at the initial stage for all conditions could be explained by the lubricating effect of the adsorbed proteins. As such, the increase in the COF would be due to damage to the mirror-like surface during the wear process. After about 8 h, the COFs in the stable stages exhibited different evolution tendencies for different electrochemical conditions: the COFs

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Fig. 13. (a) Shear direction on the contact surfaces of a cylinder on the plane in frictionally slipping contact mode; (b–d) distributing curves of the shear stress inside the plane with different depths for different applied potentials.

slightly increased at 0.2 VAg/AgCl while they deceased slowly at OCP and at −0.8 VAg/AgCl . This can be explained by the fact that the lower applied potential can increase the amount of adsorbed protein and benefit the formation of the tribo-film. Thus, the lubricating effect of the tribo-film can decrease the COF. For a high passive potential, the thickness and coverage fraction of the tribo-film all decreased and may be more easily mechanically removed from the metal surface, leading to an increase in COF. The COF and contact stress play a decisive role in the frictional shear stress. Frictional shear stress can act on the metal surface and generate stacking faults. Fig. 13a shows the shear direction on the contact surfaces of a cylinder on the plane in frictionally slipping contact mode. In the diagram, F is the applied load, v is the sliding direction, a is the semi-contact width from the Hertzian contact theory, and point (x, z) denotes a location under the surface to calculate the shear stress  xz located there. Due to the contact geometries and the loads were all the same for different applied potentials, the maximum contact pressure p0 was same and can be given by Eq. (7) from Hertzian contact theory:



p0 =

PE ∗ R

(7)

where R is the radius of the cylinder, P is the load per unit length and E* is the contact modulus. In this study, p0 = 231 MPa. The semicontact width a is given by:



a=

4PR E ∗

As such, a = 11 ␮m in this study.

(8)

The distributing curves of the shear stress could be obtained from a theoretical mechanics calculation [37]. Fig. 13b–d shows the distributing curves of the shear stress inside the plane with different depths for different applied potentials. It can be seen that the shear stresses at all depths increase as the applied potential increases, and that the maximum shear stress occurs at a depth of 0.5a. The biggest differences appear on the topmost surface (a/z = 0) with the maximum shear stress value equal to P0 ( is the COF value). The higher shear stress value might lead to severer subsurface deformation. Thus, this may be another reason why the most severe subsurface deformation appears for 0.2 VAg/AgCl . On the other hand, the tribo-film might act as a buffer layer, lowering the strain and relieving the subsurface deformation. The thickest tribo-film at −0.8 VAg/AgCl may have the best effect in relieving the subsurface deformation. However, this supposition requires further work if it is to be proven.

5. Conclusions The effect of electrochemical corrosion conditions on the subsurface microstructure evolution of a CoCrMo alloy was investigated in this study and the following conclusions can be drawn:

(1) The increase in the applied potential from cathodic to anodic can lead to a severer subsurface deformation of the CoCrMo alloy in PBS with BSA. The higher applied potential results in a thicker, severer plastic deformation layer with a higher density of stacking faults and finer grains near the top surface.

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(2) The effect of the interaction between the electrode potentials with the adsorption of proteins on the subsurface deformation was investigated. The increase in the applied potential can decrease the amount of adsorbed protein on the CoCrMo alloy surface, and this would decrease the coverage fraction and the thickness of tribo-film, and then increased the COF. These could increase the shear stresses at all depth and aggravate subsurface deformation. (3) The presence of a passive film for a high applied potential can reduce or even suppress the dislocation’s annihilation by blocking the metal surface, which also can lead to a severer subsurface deformation. Acknowledgement This work was supported by the National Natural Science Foundation of China under grant No. 51201011. References [1] A.W.E. Hodgson, S. Kurz, S. Virtanen, V. Fervel, C.O.A. Olsson, S. Mischler, Passive and transpassive behaviour of CoCrMo in simulated biological solutions, Electrochim. Acta 49 (2004) 2167–2178. [2] M.T. Mathew, M.J. Runa, M. Laurent, J.J. Jacobs, L.A. Rocha, M.A. Wimmer, Tribocorrosion behavior of CoCrMo alloy for hip prosthesis as a function of loads: a comparison between two testing systems, Wear 271 (2011) 1210–1219. [3] P. Tunthawiroon, Y. Li, A. Chiba, Influences of alloyed Si on the corrosion resistance of Co–Cr–Mo alloy to molten Al by iso-thermal oxidation in air, Corros. Sci. 100 (2015) 428–434. [4] Y. Yan, Bio-tribocorrosion in Biomaterials and Medical Implants, Elsevier, 2016. [5] Y. Sun, V. Rana, Tribocorrosion behaviour of AISI 304 stainless steel in 0.5 M NaCl solution, Mater. Chem. Phys. 129 (2011) 138–147. [6] J. Gallo, S.B. Goodman, Y.T. Konttinen, M.A. Wimmer, M. Holinka, Osteolysis around total knee arthroplasty: a review of pathogenetic mechanisms, Acta Biomater. 9 (2013) 8046–8058. [7] Y. Yan, D. Dowson, A. Neville, In-situ electrochemical study of interaction of tribology and corrosion in artificial hip prosthesis simulators, J. Mech. Behav. Biomed. Mater. 18 (2013) 191–199. [8] M.T. Clarke, P.T. Lee, A. Arora, R.N. Villar, Levels of metal ions after small- and large-diameter metal-on-metal hip arthroplasty, J. Bone Joint Surg. Br. Vol. 85 (2003) 913–917. [9] M.T. Mathew, C. Nagelli, R. Pourzal, A. Fischer, M.P. Laurent, J.J. Jacobs, M.A. Wimmer, Tribolayer formation in a metal-on-metal (MoM) hip joint: an electrochemical investigation, J. Mech. Behav. Biomed. Mater. 29 (2014) 199–212. ˜ [10] L. Casabán Julián, A. Igual Munoz, Influence of microstructure of HC CoCrMo biomedical alloys on the corrosion and wear behaviour in simulated body fluids, Tribol. Int. 44 (2011) 318–329. [11] R. Büscher, A. Fischer, Metallurgical aspects of sliding wear of fcc materials for medical applications, Materialwiss. Werkstofftech. 34 (2003) 966–975. [12] R. Pourzal, R. Theissmann, M. Morlock, A. Fischer, Micro-structural alterations within different areas of articulating surfaces of a metal-on-metal hip resurfacing system, Wear 267 (2009) 689–694. [13] Z. Wang, Y. Yan, L. Qiao, Tribocorrosion behavior of nanocrystalline metals—a review, Mater. Trans. 56 (2015) 1759–1763. [14] R. Büscher, A. Fischer, The pathways of dynamic recrystallization in all-metal hip joints, Wear 259 (2005) 887–897.

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