Materials Science and Engineering C 37 (2014) 292–304

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Tribological properties of Zr61Ti2Cu25Al12 bulk metallic glass under simulated physiological conditions Yu Wang a, Ling-ling Shi a, De-li Duan b, Shu Li b, Jian Xu a,⁎ a b

Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China Specialized Materials and Devices Division, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China

a r t i c l e

i n f o

Article history: Received 22 July 2013 Received in revised form 29 November 2013 Accepted 5 January 2014 Available online 11 January 2014 Keywords: Wear Tribocorrosion Orthopedic implant Metallic glass Zirconium alloy

a b s t r a c t In this work, wear resistance of a Zr61Ti2Cu25Al12 (ZT1) bulk metallic glass (BMG) in dry-sliding and simulated physiological media was investigated using ball-on-flat tribological approach and Si3N4 ball as counterpart. It was indicated that wear resistance of the BMG in air and deionized water is superior to Ti6Al4V alloy but inferior to 316 L stainless steel (316 L SS) and Co28Cr6Mo (CoCrMo) alloy. However, under simulated physiological media such as phosphate buffered solution (PBS) and Dulbecco's modified Eagle medium with 10 vol.% fetal bovine serum (DMEM + FBS), the ZT1 BMG exhibits decreased wear resistance in comparison with the Ti6Al4V, 316 L SS and CoCrMo. This is probably associated with its moderate pitting corrosion resistance in the medium containing chloride ions. The presence of protein in the solution has a significant effect to ruin pitting resistance of the BMG, then causing more severe wear damage. Under the dry-wear condition, abrasive wear is a predominant wear mechanism for the ZT1, whereas under deionized water, deterioration induced by abrasive wear can be mitigated. In simulated physiological media, wear deterioration is a typical tribocorrosion controlled by synergistic effects of the abrasive and corrosive wear. For the four investigated metals, wear resistance does not exhibit distinct correlation with hardness, whereas the material with high Young's modulus possesses better wear resistance. © 2014 Elsevier B.V. All rights reserved.

1. Introduction Metallic materials play a vital role in the repair or replacement of body tissue that has been diseased or damaged, including the hip, knee, finger joints and dental roots [1–3]. In comparison with ceramics and polymers, metals are more adequate as loading-bearing components, owing to their good combination of high mechanical strength and fracture toughness. Currently, representative biomedical metals that are widely used in clinics include titanium and its alloys, cobalt– chromium–molybdenum based alloys and stainless steels. Each material has its own advantages. However, wear debris/particles and release of toxic metallic ions, which are usually generated by the synergistic effects of corrosion and fretting wear, have been a long-standing challenge, which gives rise to inflammation; osteolysis related to aseptic loosening, allergy and even cancer [2,4–11]. Therefore, the development of new alloys with high corrosion and wear resistance in service under physiological environments remains considerably interesting. In contrast to the conventional crystalline metals, amorphous alloys or metallic glasses manifest substantially uniform microstructure, without the defects like dislocation and grain boundary. Periodic atomic ordering arrangements in metallic glass only occur in short-ranges rather than in long-ranges like crystalline solids. Consequently, amorphous/glassy alloys ⁎ Corresponding author. Tel.: +86 24 23971950; fax: +86 24 23971215. E-mail address: [email protected] (J. Xu). 0928-4931/$ – see front matter © 2014 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msec.2014.01.016

exhibit a number of unique properties, such as high yield strength, large elastic strain (~2%) and excellent corrosion resistance. In light of these features, zirconium-based bulk metallic glass (BMG) is recently proposed as a candidate material for prostheses implants [12–15]. On the other hand, the wear and tribology behavior of the BMGs has not been well understood so far [16–25], even in debate, since these properties are highly dependent on the service environment and counter pair, unlike conventional intrinsic mechanical properties. Fu et al. [16] showed that there is no indication of exceptional tribological properties for the Zr41.2Ti13.8Cu12.5Ni10.0Be22.5 BMG when tested with a pinon-disk geometry without lubrication. It is clear that one should not simply assume that all BMGs have exceptional tribological properties. Tam and Shek [18,19] indicated that the wear resistance of Cu-based BMG is not directly proportional to hardness and does not follow the wear law. The BMG suffers severe wear compared with the 304 stainless steel. By investigating the dry-sliding wear behavior of Zr41.2Ti15.5Cu14.5 Ni3.5Be24.5 BMG, Jin et al. [22] found that formation of the oxide layers and their subsequent peel-off constituted the main mechanism of wear. Furthermore, Chen et al. [25] showed that the wear resistance of Zr60.14Cu22.31Fe4.85Al9.7Ag3 BMG against ultra-high molecular weight polyethylene (UHMWPE) is superior to that of conventional as-cast CoCrMo alloy as counterpart under simulated physiological media. Obviously, these efforts renewed interest to investigate the tribological properties of BMGs for their potential application as biomedical implants. Recently, Espallargas et al. [26] illustrated that in the simulated

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body fluid, Zr55Cu30Ni5Al10 BMG suffered from wear accelerated corrosion generating larger wear rates due to galvanic coupling effects, in which the alumina was used as the counterpart material. More recently, a new Zr-based BMG, Zr61Ti2Cu25Al12 (designated as ZT1 hereafter), was developed [27–29]. In comparison with the previously-developed BMGs, this alloy is not only compositionally free from the toxic elements such as nickel, cobalt and beryllium, but also with unique mechanical properties including low Young's modulus (E = 83 GPa) and high fracture toughness (KJIC = 130 MPa√m). Moreover, as assessed by in vitro cellular responses [30], the ZT1 BMG manifests good biocompatibility comparable to the pure titanium (CP–Ti) and Ti–6Al–4V alloy. In addition, as reported recently [31], it is surprising to note that Zr-based particles which are used to mimic the wear debris induced less toxicity and inflammatory responses when compared with CoCrMo-alloy and Ti-alloy particles. Consequently, it is of great interest to wonder whether the ZT1 BMG is qualified to be used for the endosseous implants or hard-tissue prostheses. As is well known, currently-used materials for hard tissue prostheses deal with the metals, ceramics and polymers. Wear mechanism and degradation during service are dependent on the materials of wear couples. As the representative, the couple of femur and femur head in the artificial hip joint was assembled with different materials, such as metal-on-metal, metal-on-ceramic and metal-on-polymer [32]. In light of high hardness of the BMG, we choose a ceramic material as the counterpart in the current work for preliminary material screening. Apart from the alumina and zirconia, silicon nitride is also a candidate material as ceramic bearing [33–35]. To exclude the potential effects of the in-situ oxidized alumina or zirconia during wear process from the alloy constituents in the ZT1 BMG, silicon nitride (Si3N4) was selected as the counterpart material. In the current work, tribological properties of the ZT1 BMG sliding against the Si3N4 ball in dry and several wet media are investigated under several loading conditions. The wet media included the deionized water, phosphate buffered solution (PBS) and a conventional cell culture medium, Dulbecco's modified Eagle media plus fetal bovine serum (DMEM + FBS). Several typical metals clinically-used for medical implants such as Ti6Al4V, 316 L stainless steel (316 L SS) and Co28Cr6Mo (CoCrMo) alloys were examined in parallel as well. The wear mechanism and correlations of wear resistance with material properties such as hardness and Young's modulus are discussed. 2. Experimental As-cast Zr61Ti2Cu25Al12 (in atomic percentage) BMG plates with a dimension of 65 × 9 × 3 mm3 were fabricated by using copper mold casting of arc-melted alloy ingots [27–29]. Amorphous feature of the BMG samples was confirmed by using X-ray diffraction (XRD), taken from the cross section surface of the as-cast plates. It was performed at a Rigaku D/max 2500 diffractometer (Rigaku, Tokyo, Japan) with monochromatic Cu Kα radiation. Rectangular ZT1 plates for wear tests were mechanically ground to the size of 30.5 × 8.5 × 2.8 mm3. Three commercial materials, Ti6Al4V, 316 L SS and CoCrMo, were processed with the same dimension as that of the ZT1 BMG. All the samples were wet ground with SiC sandpapers down to 2000 grit specification, and then polished with 1.0–2.5 μm diamond paste. Final surface roughness of the samples, Ra, is less than 0.17 μm. A

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commercial ceramic Si3N4 ball with a diameter of 12.7 mm (G10 grade), supplied by Shanghai Unite Technology Co., Ltd. was used as a counterpart material. Several typical mechanical properties of the investigated materials are listed in Table 1 for comparison. A home-made wear apparatus of the ball-on-flat reciprocating sliding, coupled with coefficient of friction recorder, was used to conduct the wear process. The cycle number, tangential force, normal force and frequency were acquired simultaneously. Four sets of testing were conducted in air, deionized water, PBS and DMEM + FBS solutions. The recipe of two simulated body fluids was presented elsewhere [36,37]. Five levels of the loading were used, from 5 N to 25 N with an increment of 5 N. The testing duration, reciprocating frequency and stroke length are set as 1 h, 1 Hz and 20 mm, respectively. Under each given condition, at least three specimens were tested to ensure reproducibility. Prior to each wear test, the specimen was ultrasonically cleaned in ethanol for 5 min to eliminate the contaminants, and then dried in hot flowing air. The counterpart was also rinsed with ethanol. After testing, all the specimens were ultrasonically cleaned in acetone and ethanol for 5 min, respectively. The wear loss of each sample was determined using a balance with an accuracy of ±0.01 mg, in terms of weight change of the specimen before and after each test. Volume loss is obtained from the weight loss, converted using the relation of V = M / ρ, where the density ρ was measured by Archimedes' method. Coefficient of friction is given in a way of an average value retrieved from the whole tests. Micro-hardness of materials was measured using a Vickers microhardness tester (Qness GmbH, Austria) under a load of 200 g with a dwelling time of 15 s. Topography of worn surfaces of the specimens and wear debris was observed using LEO Supra35 scanning electron microscopy (SEM, Zeiss, Germany), attached with energy dispersive X-ray (EDX). Chemical compositions of the as-received and worn surfaces of the ZT1 BMG were analyzed by using X-ray photoelectron spectrometer (XPS). The XPS measurements were performed using an ESCALAB250 surface analysis system (Thermo VG, USA) with a monochromatic Al Kα X-ray source of 1486.6 eV. The spot size was 500 μm × 500 μm and the pass energy was 50 eV with an energy step size of 0.1 eV. The measured binding energies were calibrated referring to the C 1 s peak with the binding energy value of 284.6 eV. The curve fitting of the XPS spectra was performed with commercial XPSPEAK4.1 analysis software. Depth profile of the sample surface was determined by argon ion beam sputtering to cover an area of 2 mm × 2 mm, using the sputtering voltage of 3 kV and sample current of 2 μA. The sputtering rate of 0.1 nm/s was used to convert the sputtering time into an approximate sputter depth. Electrochemical behavior of the four investigated metals in two simulated body fluids with or without proteins was assessed by using potentiodynamic polarization tests. It was conducted on a Model 2273 electrochemical workstation (EG&G Princeton Applied Research), connected to a three electrode cell with saturated calomel as reference electrode (USCE = 241 mV) and platinum sheet as counter electrode. Before each polarization scan was initiated, the sample was kept in the electrolyte for 1 h when the variation of corrosion potential, the Ecorr, is less than 2 mV over a period of 5 min. The potentiodynamic polarization scanning with a rate of 0.17 mV/s was started at 50 mV below open circuit potential (OCP) versus USCE.

Table 1 Mechanical properties of Zr61Ti2Cu25Al12 metallic glass, 316 L stainless steel, Co28Cr6Mo, Ti6Al4V alloy and Si3N4. Material

Density (g/cm3)

Young's modulus (GPa)

Poisson ratio

Yielding strength (MPa)

Elastic strain limit (%)

Vickers hardness (GPa)

Fracture toughness (MPa√m)

Co28Cr6Mo 316 L SS Ti6Al4V Zr61Ti2Cu25Al12 Si3N4

8.27 7.93 4.38 6.43 3.23

230 210 110 83 300

0.3 0.3 0.342 0.367 0.27

897 460–690 860 1600 –

0.2 0.3 0.7 2.0 –

3.76 3.34 3.62 4.90 15.00

– 100 83 130 –

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3. Results 3.1. Tribological behavior under dry-sliding in air Fig. 1(a) shows a plot of volume loss against applied normal load in air during wear process for the ZT1 BMG and Ti6Al4V. The two alloys exhibit a similar trend which is the volume loss linearly increases with raising the applied load. Nevertheless, the variation of slopes in the two cases is quite different. Change in volume loss induced by the increased load for the Ti6Al4V is more sensitive with respect to the BMG. As the applied load is reduced down to 5 N, the difference in volume loss between the two alloys is negligible. At the load of 25 N, the volume loss of the Ti6Al4V is about 2.4 times larger than that of the ZT1. Hence, wear resistance of the ZT1 BMG under the dry-wear condition is remarkably superior to the Ti6Al4V. Our finding that the wear rate of Ti6Al4V increases with the increase of the applied load is also in agreement with the results of its dry sliding wear against hardened steel [38]. Coefficient of friction as a function of applied load for the two alloys is displayed in Fig. 1(b). As the normal load is elevated, coefficient of friction of the ZT1 BMG increases non-linearly in a small range from 0.33 at 5 N to 0.44 at 25 N, while the coefficient of friction of Ti6Al4V varies in a broad range from 0.29 to 0.58. This means that friction behavior in air for the ZT1 BMG is less sensitive to the applied load with respect to the Ti alloy. Fig. 2 illustrates the specific wear rate and coefficient of friction of the ZT1 BMG, Ti6Al4V, 316 L SS and CoCrMo, tested under the load of

Fig. 1. (a) Volume loss and (b) coefficient of friction versus normal load for Zr61Ti2Cu25Al12 BMG and Ti6Al4V alloy tested in air. The curve shown in (b) is fitted with Gaussian function.

25 N in air. The CoCrMo alloy exhibits the lowest specific wear rate, while the wear rate of 316 L SS is slightly higher than that of CoCrMo alloy. The wear rate of ZT1 BMG falls within a level between Ti6Al4V and the 316 L SS as well as CoCrMo alloy. It implies that the wear resistance of ZT1 BMG under dry-wear condition is inferior to the 316 L SS and CoCrMo, but significantly superior to Ti6Al4V. However, it is interesting to note that the difference in coefficient of friction between the four alloys is not remarkable, merely varying within 0.45–0.65. There is no distinct correlation between coefficient of friction and specific wear rate. This is not surprising since friction is a complex phenomenon that involves mechanical, chemical and physical responses from the metal and the counter piece. As indicated, the 316 L SS and CoCrMo alloy with higher coefficient of friction manifest lower specific wear rate. In contrast, the ZT1 BMG with the lowest coefficient of friction behaves as moderate wear resistance, whereas the Ti6Al4V with a coefficient of friction lower than 316 L SS and CoCrMo alloy presents the weakest wear resistance among the four investigated metals. Fig. 3(a)–(e) shows SEM images of the worn scar and collected wear debris ZT1 BMG, and surface of counterpart Si3N4 ball. As seen in Fig. 3(a), a number of ploughed grooves with width ranging from 1 to 30 μm, parallel to the sliding direction, are present in the worn surface. This is a typical topography feature of abrasive-wear controlled mechanism. The fractured junctions produced worn flakes that were trapped between the contact surfaces. The flakes were squeezed into the subsurface, and ploughed from the renewed surface under normal force. As shown in Fig. 3(b), some extruded materials in irregular shapes exist along the groove shoulder, indicating that severe plastic deformation happened locally during reciprocating sliding. The collected wear debris presented in Fig. 3(c) can be classified as two groups, according to their morphology. As observed in higher-magnification, one group is the curved chip-like foils with a length of several hundred micrometers and a width of over 500 μm, as seen in Fig. 3(d). They were produced at the initial stage of wear process, which took place in a short run-in period. During this period, contact surface of the BMG was severely cut by rough asperities of counterpart ball. Morphology of these chips is substantially similar to the machined chips of Zr-based BMG [39]. At the subsequent steady stage, flake-like debris in fragmented size is produced predominantly, which is created by reciprocating plough. Fig. 3(e) displays an appearance taken from the surface of counterpart ball. As indicated, flake-like debris in large size adhere to the surface of counterpart ball. Such large-sized smeared debris is formed through squeezing the small debris between substrate and counterpart. By virtue of EDX chemical analysis, the adhesive debris is identified as the BMG itself. To examine whether the crystallization of the BMG occurs under wear process, amorphous nature of worn surface of the ZT1 BMG is re-checked with XRD. Fig. 4 shows XRD patterns taken from the sample

Fig. 2. Specific wear rate and coefficient of friction drawn as diamond symbol for Zr61Ti2 Cu25Al12 BMG, Ti6Al4V, 316 L SS and CoCrMo alloy, tested in air under loading of 25 N.

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Fig. 3. SEM images of worn-surface topography for Zr61Ti2Cu25Al12 BMG after wear test under loading of 25 N in air, (a) lower magnification and (b) higher magnification, (c) and (d) taken from wear debris, (e) counterpart surface.

surface before and after the wear experiences. Similar to the starting BMG sample, no diffraction peaks from crystalline phase are detectable in the worn surface. This means that the wear-induced crystallization does not take place under the current conditions. It is in agreement with the previous findings of the wear performance under dry sliding for the Zr52.5Cu17.9Ni14.6Ti5Al10 [17] and Cu60Zr30Ti10 BMGs [19]. In addition, no noticeable oxidation took place either as the predominant wear mechanism, which is used to appear in some dry-sliding Zrbased BMG [22]. Fig. 5 illustrates the element concentration of surface

Fig. 4. XRD patterns taken from sample surface of Zr61Ti2Cu25Al12 BMG plate before and after wear testing.

layer for the as-received and worn BMG, which is determined by using XPS survey. With the approach to define the depth at which oxygen concentration is a half of outermost surface as the thickness of surface oxide layer, thickness of spontaneously-passive oxide film at the surface of as-cast BMG plate is estimated to be around 8 nm, as shown in Fig. 5(a). For the worn sample, the width of wear-induced groove is about 1 mm, less than the XPS beam-spot size, which ensures the collected signals are contributed only from the wear scar surface. Hence, thickness of oxide film at the worn surface is determined to be around 3 nm, as seen in Fig. 5(b). This means that thickness of repassive oxide layer under the wear processing is remarkably thinner than that of the as-received samples. Even though the thickness of the oxide film was reduced, chemical state of the film is substantially identical as the air-formed one. As depicted in Fig. 6(a) and (b), in the two cases, the Zr element detected in the surface layer is mainly present in the form of oxidized state, suggesting that the formed oxides mainly consisted of oxidized Zr 4 + species such as the ZrO 2 , Zr3(PO4)4, ZrO(OH)2 and Zr(OH)4 [40,41]. For a comparison, SEM images of worn surfaces of the Ti6Al4V, 316 L SS and CoCrMo alloy tested in air dry-sliding are illustrated in Fig. 7. As seen in Fig. 7(a), numerous grooves in a width of several-tens of micrometers appeared in the worn scars of Ti6Al4V alloy, indicative of severe wear damage. In contrast to the ZT1 BMG, the worn surface looks much rougher, together with deeper grooves. As indicated in Fig. 7(b), the grooves are partially covered at the ridge by plastically deformed layer with smeared feature, together with some pits as marked by arrows in Fig. 7(b). These evidences reveal that the Ti6Al4V alloy under dry sliding suffered abrasive wear and partial adhesive wear. Also, this behavior is in agreement with the previous findings [42]. Moreover, a number of micro-cracks perpendicular to the sliding direction are observed in the

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Fig. 6. XPS spectra taken from (a) as-received surface and (b) worn surface of Zr61Ti2Cu25 Al12 BMG. Fig. 5. Concentration of elements in surface layer determined with XPS survey versus sputtering depth for (a) as-received and (b) dry-sliding worn Zr61Ti2Cu25Al12 BMG.

worn surface, which is similar to the observation as presented in Refs. [43,44]. Fig. 7(c) illustrates the worn-scar morphology of 316 L SS, showing some grooves with the diverse widths and depths. Unlike the case of Ti6Al4V, most of the grooves are much narrower and shallower. Some patches with dark contrast, like isolated islands, exist on the worn surface. As revealed by EDX analysis, these dark patches are identified as oxides of iron, chromium and nickel elements formed on the surface, which are expected to play a role to mitigate the plunging process of the substrate. Additionally, a higher-magnification image displays that no micro-cracks as those which appeared in the Ti6Al4V are present inside the ploughed groove, as seen in Fig. 7(d). Hence, the 316 L SS just suffered slight abrasive wear in air. Fig. 7(e) shows SEM image of the worn surface of CoCrMo alloy tested in air, showing a rather smooth worn surface only with some tiny scratches. In higher-magnification observation for the local area as shown in Fig. 7(f), some very shallow pits in size around 5 μm with darker contrast reside in the surface, suggesting that only slight scratch took place in the case of the CoCrMo. 3.2. Tribological behavior under wet wear in simulated body fluids Fig. 8(a) and (b) displays the wear volume loss and coefficient of friction as a function of applied load for ZT1 BMG in two solutions, PBS and DMEM + FBS, respectively, together with the data of dry wear for comparison. As shown in Fig. 8(a), volume loss linearly increases with elevating the load in both solutions with or without proteins. This

tendency is substantially similar to the situation of dry wear in air, but the elevation of curves attained in the two corrosive media is more distinct in comparison with the curve in air. In other words, it means that wear deterioration of ZT1 BMG in simulated body fluids are more sensitive to the loading variations. As the applied load is reduced to less than 10 N, wear volume loss in the PBS without proteins is comparable to that in air. However, under larger loading such as 25 N, the wear volume loss exhibits a significant dependence of the medium. In the DMEM + FBS, wear volume loss is noticeably enhanced. Thus, a conclusion can be drawn that as a corrosive medium, simulated body fluid acts as the enhancement of the wear deterioration for ZT1 BMG. Fig. 8(b) shows a plot of coefficient of friction against applied load in air and two fluids. Different from the scenario in air, variation of coefficient of friction as a function of the applied load is not simply monotonic in the two cases. Under the loading less than 15 N, coefficient of friction in DMEM + FBS solution is greater than that in PBS solution, while under higher loading condition contrary results are presented. For comparison, Fig. 9(a) and (b) displays the specific wear rate and coefficient of friction of the four investigated metals in different media including the air, deionized water, PBS and DMEM + FBS solution, respectively. In terms of sensitivity to medium, the four materials can be categorized into two groups, as shown in Fig. 9(a). Both of the ZT1 BMG and Ti6Al4V manifest remarkable dependence of wear environment. The BMG exhibits increased wear rates in the four media in the following order: deionized water, air, PBS, and DMEM + FBS, while the wear rate of Ti6Al4V significantly increases in the following order: deionized water, PBS, DMEM + FBS and air. Thus, the wear resistance of Ti6Al4V in air and deionized water without corrosive media is inferior to that of ZT1 BMG, whereas the case is contrary in simulated body

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Fig. 7. SEM images of worn surface topography for (a) Ti6Al4V, (c) 316 L SS and (e) CoCrMo alloy tested under loading of 25 N in air. (b), (d) and (f) are zoom-in images of blocked areas in (a), (c) and (e), respectively.

fluids. The reason is likely associated with the excellent corrosion resistance of Ti6Al4V in chloride ion medium, which will be addressed later in Section 3.3. The wear properties of 316 L SS and CoCrMo alloy are quite comparative. Specific wear rate of CoCrMo alloy is slightly lower than that of 316 L SS in all media. Both alloys are less sensitive to the medium changes, unlike the ZT1 BMG and Ti6Al4V. Just due to lack of the lubrication, their wear rates in air are higher than those in aqueous solutions. Evidently, wear resistance of 316 L and CoCrMo is apparently superior to that of the ZT1 BMG and Ti6Al4V in all media. As shown in Fig. 9(b), coefficients of friction for the ZT1 in different media are very comparative, around 0.5–0.55. For the Ti6Al4V, its coefficient of friction in air is about two-fold higher than that in all three aqueous solutions, which is probably caused by the absence of lubrication. The medium effects on coefficient of friction for 316 L SS and CoCrMo are nearly similar. In both cases, coefficient of friction in media increased in the following order: DMEM + FBS, PBS, deionized water, and air. It indicates that the protein and inorganic compounds can play a role of lubrication in some extent. Fig. 10(a) and (b) illustrates SEM images of worn-surface topography for the ZT1 BMG and its counterpart Si3N4 ball after testing in deionized water, respectively. Similar to the case of dry sliding in air, numerous shallow grooves and an amount of entrapped wear debris in flake shape are present in the worn surface, as seen in Fig. 10(a). Because of the presence of liquid phase between counter surfaces, peel-off debris is easily removed away from the contact area. In contrast, surface of the counterpart ball is clean and smooth, as shown in Fig. 10(b), without visible trace of material transfer. This suggests that adherence of wear debris to the surface is difficult due to flush effect of liquid phase.

Fig. 11(a)–(d) displays SEM images of worn-surface topography of Ti6Al4V, 316 L SS and CoCrMo tested in deionized water. As shown in Fig. 11(a), discontinuous grooves and traces of delamination are observed in the surface of Ti6Al4V. With lubrication of non-corrosive liquid phase, plough-induced damage is significantly mitigated in comparison with dry sliding in air [seen in Fig. 7(a)], but severe plastic deformation remains operative during wear process. Amounts of extruded materials are spread out along the groove sides. In highermagnification observation as shown in Fig. 11(b), numerous microcracks perpendicular to sliding direction exist in the worn surface, similar to dry-wear scenario [seen in Fig. 7(b)]. Therefore, wear mechanism of Ti6Al4V under wet sliding condition is basically similar to the case of dry sliding, which is predominated by abrasive wear. Fig. 11(c) shows the worn-surface topography of 316 L SS. In contrast to dry-wear scenario [seen in Fig. 7(c) and (d)], ploughed deep grooves disappeared, and only the traces of slight scratching are present. Uniform oxide film formed on the surface of 316 L SS bonds well to the substrate, and plays a role of protecting the subsurface [45]. As shown in Fig. 11(d), topography of worn-surface of CoCrMo is substantially comparative to the dry-wear case [seen in Fig. 7(e)], indicating that predominant wear mechanism remains as slight abrasion. Fig. 12(a)–(f) illustrates SEM images of worn-surface topography of the four investigated metals tested in PBS. As seen in Fig. 12(a) and (b), wear scars of ZT1 BMG are substantially comparable to those tested in deionized water [seen in Fig. 10(a) and (b)]. Small amount of the BMG is transferred as it adheres to the counterpart surface (not shown here), indicating that the wear debris are not so easily removed away in PBS as it does in deionized water. This is likely due to that the PBS medium is more viscous than the water. As observed, most of the wear

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Fig. 9. (a) Specific wear rate and (b) coefficient of friction for Zr61Ti2Cu25Al12, Ti6Al4V, 316 L stainless steel and CoCrMo alloys, tested in air, deionized water, PBS and DMEM + FBS media under loading of 25 N.

Fig. 8. (a) Volume loss and (b) coefficient of friction for Zr61Ti2Cu25Al12 BMG in PBS and DMEM + FBS media. Data of dry wear are given as well for comparison.

debris keeps flake-like shape. As shown in Fig. 12(c), topography of Ti6Al4V is similar to the case in deionized water [seen in Fig. 11(a)]. Numerous micro-cracks appear in the worn surfaces again. This happened in all the currently-used media. Furthermore, these findings are in accord with a fact that the worn damage of Ti6Al4V alloy was conducted in Hank's solution [46]. Thus, it can be concluded that microcracks formed on the worn surface of a Ti6Al4V are independent on operating medium, which are intrinsically related to its weak resistance to shear deformation. Fig. 12(e) and (f) displays worn-surface topography of the 316 L SS and CoCrMo alloy after wear processing in the PBS, respectively. Only slight scratches are present on smooth surfaces in both cases, showing that both of 316 L SS and CoCrMo are subjected to only slight abrasive wear. No distinct difference exists in wear mechanism, irrespective of the wear media with or without chloride ions. Fig. 13 shows SEM images of worn-surface topography of the four investigated metals tested in the protein-containing DMEM + FBS solution. As seen in Fig. 13(a), worn surface of the ZT1 BMG presents typical abrasive wear feature similar to that in PBS [seen in Fig. 12(a) and (b)], although the specific wear rate increases by about 23%. Shallow grooves and small amount of wear debris are spread on the worn surface. Morphology of removed wear debris is nearly the same as that collected in PBS without protein. As seen in Fig. 13(b), similar to the case tested in deionized water and PBS, severe abrasion appeared in the worn surface of Ti6Al4V. For the 316 L SS, ploughed grooves are evident in the surface, as seen in Fig. 13(c), indicating more severe wear damage compared with those in PBS. For the CoCrMo, as seen in Fig. 13(d), only slight abrasive wear happens, reflecting its excellent

wear resistance. Its wear mechanism seems irrespective of the nature of aqueous solution, such as the presence of inorganic ions, molecules and protein.

Fig. 10. SEM images of worn-surface topography for (a) Zr61Ti2Cu25Al12 BMG and (b) Si3N4 counterpart after wear in deionized water tested under loading 25 N.

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Fig. 11. SEM images of worn-surface topography for (a) Ti6Al4V, (c) 316 LSS and (d) CoCrMo alloy in deionized water tested under loading of 25 N. (b) is zoom-in image of blocked area in (a).

3.3. Potentiodynamic polarization Fig. 14(a) and (b) displays potentiodynamic polarization curves of the four investigated metals in PBS, and curves of the ZT1 BMG and Ti6Al4V in DMEM + FBS, respectively. It is demonstrated that all materials are passivated before pitting occurs, together with low passive

current densities in passivation region. From the polarization curves, using an extrapolative method of Tafel line, we measured corrosioncurrent density (icorr). According to Faraday's law, corrosion penetration rate (CPR, μm/yr) was calculated: CPR = 0.327 Micorr / mρ, where M (g/ mol), m, and ρ (g/cm3) are the atomic-fraction-weighted values of atomic weight, ion valence, and density, respectively, for the alloy

Fig. 12. SEM images of worn-surface topography for (a) Zr61Ti2Cu25Al12 BMG, (c) Ti6Al4V, (e) 316 L SS and (f) CoCrMo in PBS tested under loading of 25 N. (b) and (d) are zoom-in images of blocked areas in (a) and (c), respectively.

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Fig. 13. SEM images of worn-surface topography for (a) Zr61Ti2Cu25Al12 BMG, (b) Ti6Al4V, (c) 316 L SS and (d) CoCrMo alloy tested in DMEM + FBS under loading of 25 N.

elements, and icorr (mA/m2) is the corrosion-current density [47]. The measured electrochemical parameters are listed in Table 2. In the PBS electrolyte, the ZT1 BMG presents the passivation at a lower pitting overpotential, as ΔE (Epit − Ecorr) = 953 mV. The mean

CPR was calculated as 0.23 μm/y. From the curve of 316 L SS, pitting overpotential is measured to be 1017 mV, which is statistically comparable to that of the ZT1 BMG. Nevertheless, 316 L SS possesses a relatively higher pitting potential (Epit), reflecting a pitting resistance better than the ZT1 BMG. In contrast, no potential plateau appears in the curves of Ti6Al4V and CoCrMo during testing process, indicating their excellent resistance to localized corrosion. The mean CPR value of the Ti6Al4V, CoCrMo and 316 L SS is determined to be 0.28, 0.49 and 1.60 μm/y, respectively. These findings are approximately in agreement with the previous results reported by Morrison et al. [37]. In their work, Zr52.5Cu17.9Ni14.6Al10Ti5 BMG (Vit105) and 316 L SS are susceptible to localized corrosion in the PBS medium, in contrast to Ti6Al4V and CoCrMo alloys. Using the CPR as a measure, the resistance to uniform corrosion for Vit105 BMG is statistically comparable to the Ti6Al4V and CoCrMo, and superior to 316 L SS. As shown in Fig. 14(b), the curves of ZT1 and Ti6Al4V show a trend similar to those in PBS, apart from that the Ecorr value shifts to the more negative side. Compared with the case in PBS, value of Ecorr, ΔE and CPR for the ZT1 BMG is −676 mV, 843 mV and 0.59 μm/y, respectively. This suggests that pitting corrosion of ZT1 BMG happens more easily in protein-containing medium of DMEM + FBS. In other words, the resistance either to the localized or to uniform corrosion is weakened due to the presence of protein. However, no potential plateau is present overall the tested potential for the Ti6Al4V, showing excellent resistance to pitting corrosion even in the protein-containing medium. The Ecorr is somewhat reduced, which is accompanied by a subtle increase of the CPR. This implies that the simulated body fluid containing protein, such as DMEM + FBS, has a severer corrosive effect than PBS does. 4. Discussion 4.1. Effect of corrosive nature of aqueous-solution on wear deterioration

Fig. 14. Potentiodynamic polarization curves of (a) Zr61Ti2Cu25Al12 BMG, Ti6Al4V, 316 L SS and CoCrMo alloy in PBS at 37 °C, (b) Zr61Ti2Cu25Al12 BMG and Ti6Al4V in DMEM + FBS solution.

According to the Archard's equation [48], wear volume loss is expected to increase linearly with elevating the normal load. In the current work, wear behavior of ZT1 BMG follows Archard's equation very well both in air and in two simulated body fluids, as seen in Fig. 8(a). It is also in accordance with several previous reports [23,49]. However, it is noteworthy that the correlation of coefficient of friction with normal load is significantly different from findings in Refs. [16,23,49]. It suggests that the relationship between coefficient of friction and applied load is dependent on the counter pair.

Y. Wang et al. / Materials Science and Engineering C 37 (2014) 292–304

301

Table 2 Electrochemical parameters related to corrosion resistance for the four investigated metals. Material

Solution

Ecorr (mV)

Zr61Ti2Cu25Al12 Ti6Al4V 316 L SS CoCrMo Zr61Ti2Cu25Al12 Ti6Al4V

PBS PBS PBS PBS DMEM + FBS DMEM + FBS

−551 −571 −234 −320 −676 −599

± ± ± ± ± ±

17 11 13 30 25 45

Epit − Ecorr (mV)

Icorr (μA/cm2)

CPR (μm/y)

953 ± 205 – 1017 ± 186 – 843 ± 22 –

0.02 0.03 0.15 0.06 0.05 0.04

0.23 0.28 1.60 0.49 0.59 0.36

As indicated in Section 3.2, wear behavior of the ZT1 in aqueous solutions exhibits a significant dependence on chemical nature of the media. Wear-induced deterioration is promoted in the following sequence: deionized water, air, PBS and DMEM + FBS. The noncorrosive liquid medium seems to act as the boundary lubrication film, then to mitigate the wear damage. Consequently, wear deterioration of the ZT1 BMG in deionized water without corrosive agent is mitigated in contrast to the case of dry wear. As observed in SEM, worn surface of the ZT1 BMG tested in PBS is much smoother than that in air, comparative to the case in deionized water. It indicates that the contact surface between ZT1 BMG and counterpart is also lubricated by PBS. However, wear deterioration of the ZT1 BMG in this medium is enhanced in comparison with the deionized water and air. From the viewpoint of electrochemical corrosion, concentration of chloride ion in the medium is a key factor responsible for corrosion of Zr-based BMG in the PBS. It is well recognized that chloride ion possesses the power to ruin the surface oxide layer, then to induce the pitting corrosion [37,50]. Gebert et al. [51,52] suggested that the chemical and physical defects in the as-cast BMG samples were preferential sites for pitting initiation, which leads to the pit growth and propagation. Selective dissolution of elements Zr and Al as well as other valve components gives rise to an enrichment of Cu element in the pit zone. The accumulated Cu element was expected to locally interact with chloride ion and to form cuprous chloride CuCl, which subsequently induces the formation of cuprous oxide Cu2O. The local accumulation of Cu-rich species may generate the galvanic coupling effects to trigger local dissolution. In the current work, it is displayed in Fig. 14(a) that the pitting corrosion of ZT1 BMG occurs at the potential over 400 mV, which is nearly half of 316 L SS, while no pitting corrosion took place in the cases of both Ti6Al4V and CoCrMo, as presented in their polarization curves. It reflects that the pitting resistance of ZT1 BMG in simulated body fluid is inferior to that of 316 L SS, Ti6Al4V and CoCrMo. It seems very consistent with the wear resistance of ZT1 BMG in simulated physiological media, which is poorer than the other three alloys. Evidently, improvement of pitting corrosion resistance in physiological media for the Zr-based BMG is critical issue in future work. 4.2. Effects of protein in simulated body fluids Under the condition of DMEM + FBS solution, tribological property of the four investigated metals is remarkably different from the case in PBS without protein, showing a general trend that coefficients of friction of 316 L SS and CoCrMo are reduced, whereas the specific wear rate of ZT1 BMG, Ti6Al4V and 316 L SS is increased, as shown in Fig. 9. Such differences in the wear and friction behavior are mainly attributed to the protein effects in simulated body fluids, which are adsorbed on the surfaces of specimens. Effects of protein on tribological property of bio-implant metals have been considerably of concern [53–57]. Scholes et al. [58] found that the effect of bovine serum albumin (BSA) on friction behavior was dependent on the counter pairs. For the metal-on-metal counter pair, the frictional factor is reduced in 100% BSA solution compared to that in carboxy methyl cellulose solution, whereas the opposite propensity was present in the case of the metal-on-UHMWPE pair or ceramic-onceramic counter pairs. It is believed that the protein adsorbed on the

± ± ± ± ± ±

0.01 0.01 0.03 0.04 0.01 0.01

± ± ± ± ± ±

Ip (μA/cm2) 0.11 0.07 0.30 0.31 0.08 0.11

0.85 0.96 0.83 0.42 1.25 1.24

± ± ± ± ± ±

0.16 0.21 0.03 0.19 0.09 0.13

surfaces of counter pairs extends the lubrication regime, resulting in a well lubricating effect for metal-on-metal counter pair. As noted by Gispert et al. [54], for the counter pairs of CoCrMo/UHMWPE and 316 L/UHMWPE, coefficient of friction is markedly reduced. In BSAcontaining Hank's solution, friction process became more stable compared to BSA-free Hank's solution. However, the BSA does not play a significant role to protect the surfaces of alumina/UHMWPE. In our work, coefficient of friction of 316 L SS/Si3N4 and CoCrMo/Si3N4 counter pairs is remarkably reduced about 50% and 60% in DMEM + FBS solution with respect to the case of protein-free PBS. But for the ZT1/ Si3N4 and Ti6Al4V/Si3N4 counter pairs, no distinct difference is found in the solutions with or without proteins. This is probably caused by a fact that severe wear damage ruins the adhesion of protein layer on the surface, then reducing the lubrication effect. The protein effect on wear properties in simulated body fluid for biomedical metals is very complex. It is far from well understanding so far. Gispert et al. [54] reported that the wear rate of UHMWPE pair against TiN-coated stainless steel was significantly mitigated in Hank's solution with bovine serum. It is associated with the effect that adsorbed protein layer suppresses the interaction of contact surfaces. Amaral et al. [59] indicate, however, that the uniform lubricating film can be destroyed by adsorbed protein layer, resulting in a large stress in certain contact area then giving rise to more severe wear. In our work, the specific wear rates of ZT1 BMG and Ti6Al4V significantly increase in DMEM + FBS solution, compared with the case in PBS. On the other hand, as indicated by potentiodynamic polarization (see Section 3.3), corrosion in DMEM + FBS for ZT1 BMG and Ti6Al4V is enhanced. In other words, the effect of protein to promote corrosion is a plausible factor responsible for the enhanced wear damage. 4.3. Correlation of wear resistance with hardness and Young's modulus of materials According to Archard's equation [48], volume loss of alloy exhibits an inverse proportional relationship with its hardness, which intuitively suggests that a material with higher hardness is expected to possess better wear resistance. It was indicated in Ref. [60] that the higher hardness of alloy corresponds to a lower wear rate. The data collected from metallic glasses and crystalline alloys were plotted together. However, no well-defined correlation is present [61]. In our work, we plot the data of specific wear rate in various media against Vickers' hardness for the four investigated metals, as shown in Fig. 15(a). It does not display a distinct correlation between the two properties. The “softest” 316 L SS among the four materials manifests wear resistance better than Ti6Al4V and ZT1 BMG. On the contrary, the ZT1 BMG with higher hardness shows poor wear resistance, especially in simulated physiological media. Consequently, it is worthy to be emphasized that it is inadequate to rank the wear resistance of alloys simply based on their hardness values. As reviewed by Greer et al. [62], pure metals showed higher wear resistance in relation to their lower hardness, which is due to their work-hardening ability. For hardened alloys and amorphous alloys, wear resistance was inferior to that of pure metals with same level hardness because of their absence in work-hardening ability. In the case of ceramics, their brittle nature led to the poorest wear resistance although their hardness is the highest. In addition, service

302

Y. Wang et al. / Materials Science and Engineering C 37 (2014) 292–304 Table 3 The calculated contact mechanics parameters of current counterparts according to the Hertzian contact mechanics, including the contact radius a, apparent contact area A, maximum pressure pmax and mean pressure pm. Counterpart

L (N)

Rball (mm)

E⁎ (GPa)

a (mm)

A (mm2)

pmax (MPa)

pm (MPa)

Si3N4–CoCrMo Si3N4–316 L SS Si3N4–Ti6Al4V Si3N4–ZT1

25 25 25 25

6.3 6.3 6.3 6.3

142 135 90 74

0.0940 0.0956 0.1095 0.1169

0.028 0.029 0.038 0.043

1350 1305 996 874

900 870 664 583

2

A ¼ πa ¼ π

Fig. 15. Plot of specific wear rate tested in different media under loading of 25 N against (a) Vickers micro-hardness and (b) Young's modulus.

environments remarkably influence the wear performance of some alloys. As seen in Fig. 9(a), wear property of Ti6Al4V and ZT1 BMG is very sensitive to service media. In the solutions with chloride ion as well as proteins, the wear damage exhibits a large variation. Hence, it is inadequate to judge the wear resistance of alloys only in terms of their hardness. Moreover, Fig. 15(b) shows a plot of wear rate against Young's modulus for the four investigated metals. It displays a general trend that the alloy with high Young's modulus behaves as a lower specific wear rate. Such a trend is the case even in simulated body fluids. As a matter of fact, Greer et al. [62] suggested a good correlation between wear resistance and Young's modulus existed only concerning the family of amorphous alloys. It is well known that in tribological contacts, the real contact area between two rubbing bodies normally differs from the nominal (geometrical) contact area. In order to address the correlation of wear rate with Young's modulus, the apparent contact area and stress on the wear surface were estimated using the Hertz contact theory [63]. According to the Hertzian contact mechanics, the contact radius a, the apparent contact area A, and both the maximum pressure pmax and the mean pressure pm can be calculated using Eqs. (1)–(5), and the data are summarized in Table 3.

a¼




1 3LR =3 4E

E ¼

1−ν 21 1−ν 22 þ E1 E2

ð1Þ

!−1 ð2Þ

ð3Þ

3L 2πa2

ð4Þ

2 p 3 max

ð5Þ

p max ¼

pm ¼


2 3LR =3 4E

where the L is the applied normal load, R is the radius of Si3N4 ball, E* is the combined Young's modulus, ν1 and ν2 are the Poisson ratios of the Si3N4 ball and tested material, respectively, and E1 and E2 are the Young's moduli of the Si3N4 ball and tested material, respectively. The data in the Table 3 show a descending trend for apparent contact area, A, as an order of CoCrMo, 316 L SS, Ti6Al4 and ZT1 with the reduction of Young's modulus. It is approximately consistent with the ranking of wear rate for the four tested materials besides the case of dry wear in air. The apparent contact area of 316 L SS and CoCrMo is comparative and about 30% less than that of Ti6Al4 and ZT1. As above-mentioned, the abrasive wear is a predominant wear mechanism for the CoCrMo, 316 L SS and ZT1 alloy, either in dry sliding or in simulated physiological solutions. In our work, the asperities of Si3N4 ball were indented into and ploughed the surfaces of tested materials, causing wear particles to be dropped from the substrates of tested materials. Therefore, by expanding the apparent contact area, more effective asperities for the tested materials plugging are involved, resulting in the acceleration of their wear rate. For the Ti6Al4V worn in simulated physiological solutions, its worn surface shows typical characteristics of abrasive wear, as seen in Figs. 12(b) and 13(b). It is associated with its less apparent contact area and excellent pitting resistance, responsible for its wear rate lower than that of the ZT1. However, in the case of dry wear, topography of the worn Ti6Al4V manifests severer grooves ploughed by asperities of counter ball and the more pits left by material delamination. This indicates that the wear mechanism consists of severe abrasive wear and adhesive wear. Synergistic effect of abrasive wear and adhesive wear accelerates its wear rate, resulting in poor wear resistance with respect to the ZT1. The analysis using the Hertzian contact mechanics indicates that the elastic contacts between the current counterparts are predominant. Consequently, it is not surprising that the higher Young's modulus of materials scales with better wear resistance, as seen in Fig. 15(b). In fact, wear damage of a material in simulated biological environments has been widely treated as a tribocorrosion process [53,64–67]. Tribocorrosion is a complex chemical–electrochemical–mechanical process leading to degradation of materials in tribological contact immersed in a corrosive environment. Synergistic effects of the combined action of corrosion and mechanical loading can accentuate the wearcorrosion rate. During this process, depending on its chemical composition and mechanical properties, passivating film may protect the surface against wear or lead to increase degradation. In this sense, depassivation–repassivation kinetics induced by the passivating film breakdown plays an important role [68,69]. This was supported by a finding that the oxide film formed on Ti6Al4V suffers severe wear

Y. Wang et al. / Materials Science and Engineering C 37 (2014) 292–304

with respect to the 316 L SS [68]. So far, few studies on the repassivation of Zr-based BMG are reported [70,71]. Peter et al. [70] roughly investigated the repassivation behavior of the Zr52.5Cu17.9Ni14.6Ti5Al10 BMG in the 0.05 M Na2SO4 solution (essentially water). It was found that when the protective passive film was removed by scratching with a diamond stylus, it was found to quickly reform. This suggested that the BMG has the ability for rapid repassivation. As a result, the alternative processes of the depassivation and repassivation for passivating film of the Zr-based BMG, similar to the Ti6Al4V, are probably a major source responsible for the enhanced wear damage in simulated biological medium. Finally, it is worthy to emphasize that even though the current Zr61 Ti2Cu25Al12 BMG does not present an advantage in wear resistance under our limited testing conditions, it should not preclude the performance of Zr-based BMGs in other types of tribosystems, for example, to use the pure Ti, alumina and UHMWPE as counterpart. In addition, there exists a wide composition range of BMGs to be explored. Furthermore, corrosion and wear resistance of the BMG are rather promising to be improved by composition optimization and surface modification, such as ion implanting and coating. 5. Conclusions and outlook With Si3N4 as wear counterpart, it is revealed that the wear resistance of Zr61Ti2Cu25Al12 bulk metallic glass in air and deionized water is superior to Ti6Al4V alloy but inferior to 316 L stainless steel and CoCrMo alloy. However, under simulated physiological media such as PBS and DMEM + FBS, the BMG exhibits decreased wear resistance in comparison with Ti6Al4V, 316 L SS and CoCrMo. This is probably associated with its moderate pitting corrosion resistance and depassivation– repassivation kinetics induced by the passivating film breakdown in the chloride-ion containing solution. The presence of protein in solution has a significant effect to ruin pitting resistance of the Zr-based BMG, causing more severe wear damage. Therefore, improvement of pitting resistance in physiological media for the Zr-based BMG is critical issue for its application as biomedical implants. In addition, screening a right material as the counterpart to couple with Zr-based BMG is an additional key factor, to ensure its lower wear rate. For the ZT1 BMG under the dry-wear condition, abrasive wear is a predominant wear mechanism. While under the cooling and lubricating conditions with deionized water, deterioration caused by abrasive wear can be mitigated. In simulated physiological environment with corrosive chloride ion, the wear process is a typical tribocorrosion approach controlled by synergistic effects of abrasive and corrosive wear. For the investigated bio-metals and Zr-based BMG, wear resistance does not exhibit distinct correlation with hardness, whereas the wear resistance is related to the Young's modulus of a material, which can be understood on the basis of Hertzian contact mechanics. Acknowledgments The authors gratefully acknowledge the assistance in wear performance tests from Mr. Zi-Yang Hu, Prof. K. Yang for providing the 316 L stainless steel samples and Prof. Q.J. Wang for providing the Ti–6Al–4V alloy samples. This work was supported through the SYNL by the National Natural Science Foundation of China under Grant No. 51171180 and No. 51001099. References [1] [2] [3] [4] [5]

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