Wear response of heat-treated Ti–13Zr–13Nb alloy in dry condition and simulated body fluid

Wear response of heat-treated Ti–13Zr–13Nb alloy in dry condition and simulated body fluid

Available online at www.sciencedirect.com Wear 264 (2008) 1015–1025 Wear response of heat-treated Ti–13Zr–13Nb alloy in dry condition and simulated ...

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Available online at www.sciencedirect.com

Wear 264 (2008) 1015–1025

Wear response of heat-treated Ti–13Zr–13Nb alloy in dry condition and simulated body fluid P. Majumdar ∗ , S.B. Singh, M. Chakraborty Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Kharagpur, Kharagpur 721302, India Received 24 August 2006; received in revised form 8 June 2007; accepted 16 August 2007 Available online 29 September 2007

Abstract The effect of heat treatment on the microstructure, hardness and sliding wear behavior of Ti–13Zr–13Nb has been investigated. The wear behavior was studied in dry as well as in wet conditions using Hank’s solution and bovine serum as the lubricating media. The microstructure of the heat-treated alloy consisted of acicular and globular ␣, ␤ and martensite. The hardness was found to depend on the volume fraction and distribution of ␣ phase in the matrix. In dry condition, there was no significant variation of the wear rate with the heat-treatment conditions except for the sample deformed above the ␤ transition temperature and aged after solution treatment at 800 ◦ C. In Hank’s solution, the wear rate of all the samples increased considerably. The rate of wear was the highest in bovine serum, possibly due to the presence of protein in the lubricant. The wear process is characterized by ridged wear scars parallel to the sliding direction with superficial plastic deformation along the direction of the sliding and smeared wear surface. The wear mechanism is mainly abrasive. © 2007 Elsevier B.V. All rights reserved. Keywords: Titanium alloy; Microstructure; Hardness; Friction and wear

1. Introduction Titanium and its alloys have become one of the most attractive classes of biomedical implant materials. They are generally preferred to stainless steels and Co–Cr alloys because of their low density, superior biocompatibility and corrosion resistance, good mechanical properties and low elastic modulus. Wear of the implant materials is a significant clinical problem [1–3]. The wear of prosthetic components generates debris, which cause adverse cellular response and inflammation, eventually leading to implant loosening and pain in the body. The presence of particulate corrosion and wear products in the tissues surrounding the implant may result in a cascade of events leading to bone loss [3–6]. The poor tribological properties of titanium and its alloys prevent their applications as tribological components [7]. The first generation orthopedic ␣ + ␤ titanium alloys such as Ti–6Al–4V ELI (extra low interstitial), Ti–6Al–7Nb and Ti–5Al–2.5Fe are already in use. In recent years, second-



Corresponding author. Tel.: +91 3222 283290; fax: +91 3222 282280. E-mail addresses: [email protected], [email protected] (P. Majumdar). 0043-1648/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.wear.2007.08.005

generation low modulus near ␤ and ␤ type titanium alloys have been developed for orthopedic applications in order to avoid the ‘stress shielding’ effect caused by the modulus mismatch between the implant and the bone. However, the wear resistance of titanium alloys is inferior when compared with Co–Cr alloys and hence, wear debris are often found in the tissues near Ti/Tialloy implants [5,8]. The common approaches used to improve the wear resistance of materials include surface modification, composition adjustment and heat treatment [3,5,7,9–11]. A fundamental understanding of the friction and wear mechanisms of titanium and its alloys is required in order to achieve enhanced performance in terms of the wear behavior. Also, proper evaluation of the wear characteristics of biomedical titanium alloys in simulated body fluid is essential. This is done using laboratory wear tests. These laboratory tests do not exactly reproduce the actual conditions in the human joint which is rather complicated and not completely understood [8,12]. Nevertheless, they are useful for predicting the wear behavior and performance of the material in vivo and are widely used for first level evaluation of the material and characterization of the wear behavior [12–21]. Such tests are also valuable for comparison of different materials, heat-treatment and surface-treatment procedures. The conditions for laboratory wear tests of implant

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materials have been spelt out by Dumbleton [12]. Both dry and wear tests have been described. Though the dry wear tests do not replicate the actual conditions in the human body, they are performed to highlight the effect of the lubricants. A number of lubricating liquids has been listed for wet tests; these include distilled water, saline, Ringers or Hank’s solutions, human or animal serum, animal synovial fluid, etc. [5,12]. Lubricants like Hank’s solution form a thin layer on the articulating surfaces and prevent direct contact between them and thereby affect the friction and wear behavior. A large body of literature exists where one or more of these lubricants have been used for evaluating the wear characteristics of implant materials [5,13–21]. For example Hank’s solution was used for tests reported in [13–21], bovine serum in [14] and a mixture of Hank’s solution and bovine serum in [13,14,17]. Accordingly, the effect of heat treatment on the wear behavior of Ti–13Zr–13Nb alloy in dry condition, in Hank’s solution and in bovine serum was studied in the present investigation. 2. Experimental Ti–13Zr–13Nb (TZN) alloy (compositions in wt%) was prepared by arc melting with a non-consumable tungsten electrode in a vacuum arc melting unit supplied by Vacuum Techniques Pvt. Ltd., Bangalore. The melting chamber was first evacuated to less than 5 × 10−6 mbar and then flushed with high-purity argon and evacuated again; the chamber was finally back-filled with the same gas before melting. The ingots were turned over and remelted at least six times in order to attain homogeneity in composition and microstructure. Three or four small buttons were remelted together and cast into ingot (approximately 40 g) in the same melting unit. The ingot was given 30–40% reduction by groove rolling at two different temperatures (800 and 650 ◦ C) and then air-cooled to room temperature. The hot rolling temperatures were selected in such a way that it was above the ␤ transition temperature in one case (800 ◦ C) and below that temperature in another (650 ◦ C). The hot rolled samples were solution treated at 800, 700 and 650 ◦ C for 1 h in dynamic argon atmosphere; this was followed by furnace cooling (FC), air cooling (AC) or water quenching (WQ). Water-quenched samples were aged at 500 ◦ C for 5 h. Room temperature X-ray diffraction analysis was carried out in a Philips, Holland, PW 1710 X-ray Diffractometer with Cu K␣ radiation at 40 kV and 20 mA. The scanning rate was kept at 3◦ 2θ/min, from 2θ = 20◦ to 100◦ . Optical microscopy was used for microstructural examination of the heat-treated samples. The polished samples were etched with Kroll’s reagent (10 vol% HF and 5 vol% HNO3 in water). Hardness of the pre-polished samples was measured at 10 kg load using a Vickers hardness tester (VM 50, Fuel Instruments & Engineers Pvt. Ltd., India); an average of 10 readings was taken. Sliding wear test was started at ambient temperature (30 ± 2 ◦ C) using a pin-on-disc wear-testing machine supplied by Ducom, Bangalore. The disc was made of hardened high chromium steel having a hardness of HRc 69. The wear samples were prepared by machining 8 mm diameter pins from

rolled bars, which were then heat treated as described above. Prior to the wear tests, the pin and the disc were polished metallographically to a roughness (Ra ) level of 0.15–0.17 and 0.4–0.55 ␮m, respectively. Tests were carried out in dry condition as well as in wet conditions. Wet wear tests were carried out with Hank’s solution and with bovine serum. As discussed earlier, both these lubricating agents are commonly employed as model fluids for simulating human body fluid conditions [5,12–21]. Freshly prepared Hank’s solution was used for each experiment in the present investigation; it had the following chemical composition (values in g/l): 8 g NaCl, 1 g glucose, 0.4 g KCl, 0.35 g NaHCO3 , 0.14 g CaCl2 , 0.1 g MgCl2 ·6H2 O, 0.06 g Na2 HPO4 ·2H2 O, 0.06 g KH2 PO4 and 0.06 g MgSO4 ·7H2 O. The pH of the solution was maintained at 7.4. Loads acting on human joints vary considerably from joint to joint. For a particular joint, it varies with time during the loading cycle (e.g., from “heel strike” to “toe off” during walking) and with loading rate (e.g., with pace (slow, normal or fast) of walking) [12,22]. It has been reported that the stresses in the living joints are of the order of 1 MPa [12]. Gispert et al. [17] used a normal pressure of 0.88 MPa in the contact zone to study the wear behavior of prosthetic TiN-coated stainless steel/ultrahigh molecular weight polyethylene (UHMWPE) tribological pair. Consistent with these reports, a normal load of 50 N (equivalent to a nominal contact pressure of 1 MPa) was applied to the specimen. A linear sliding speed of 1 m/s was maintained. Tests were continued up to a sliding distance of 2000 m; this was sufficient for attainment of steady-state situation in all the cases. Wear rate, in the steady state, was computed by the rate of height reduction of the pin. The height reduction data were converted into volume loss by multiplying it with the area of cross-section of the test pin. Wear surfaces, wear debris and subsurface were examined using scanning electron microscopy (SEM) and optical microscopy. 3. Results and discussions The ␤ transition temperature of Ti–13Zr–13Nb has been reported to be 735 ◦ C [4]. The heat-treatment temperatures in the present study were selected based on this. The microstructures of the heat-treated samples were found to be consistent with those reported in literature in similar cases and is described below. 3.1. Microstructure The microstructures of the samples deformed at 800 and 650 ◦ C (above and below the ␤ transition temperature, respectively) and followed by different sets of heat-treatment conditions are shown in Figs. 1–4. Samples deformed above the ␤ transus, i.e. at 800 ◦ C and solution treated at 800 ◦ C followed by furnace cooling showed a basket-wave structure (Fig. 1Ia) formed from prior ␤ grains. Fine ␣–␤ structure was observed within prior ␤ grains in the air-cooled sample (Fig. 1Ib). Presence of primary ␣ and retained ␤ on a fine scale was observed for samples solution treated at 650 ◦ C (Fig. 1IIa–IIc). Water quenching from 800 ◦ C resulted in formation of martensite with some amount of retained ␤ phase (Fig. 1Ic). Martensite was not

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Fig. 1. Microstructures of the Ti–13Zr–13Nb alloy deformed at 800 ◦ C and solution treated at (I) 800 ◦ C and (II) 650 ◦ C for 1 h followed by (a) FC, (b) AC and (c) WQ.

Fig. 2. Microstructures of the samples aged at 500 ◦ C for 5 h after water quenched from (a) 800 ◦ C and (b) 650 ◦ C.

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Fig. 3. Microstructures of the Ti–13Zr–13Nb alloy deformed at 650 ◦ C and solution treated at (I) 700 ◦ C and (II) 650 ◦ C for 1 h followed by (a) FC, (b) AC and (c) WQ.

Fig. 4. Microstructures of the samples aged at 500 ◦ C for 5 h after water quenched from (a) 700 ◦ C and (b) 650 ◦ C.

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Fig. 5. XRD pattern of homogenized Ti–13Zr–13Nb alloy deformed at 800 ◦ C and solution treated at (a) 800 ◦ C and (b) 650 ◦ C for 1 h followed by FC, AC and WQ.

observed in sample water quenched from 650 ◦ C. Partitioning of alloying elements takes place during ␤ to ␣ transformation. The ␤ phase is enriched with Nb during solution treatment at 650 ◦ C. This reduces the Ms temperature of the untransformed ␤ to below room temperature [23] and thus no martensite was formed on water quenching from 650 ◦ C. Aging at 500 ◦ C for 5 h after quenching from a solution treatment temperature of 800 ◦ C transformed the martensite into ␣ and ␤ phases and produced a uniform distribution of acicular and globular ␣ (Fig. 2a). A similar aging of the sample water quenched from the solution treatment temperature of 650 ◦ C resulted in the growth of globular ␣ (Fig. 2b). For the samples deformed below the ␤ transus, i.e. at 650 ◦ C, the overall microstructure (Fig. 3) after heat treatment at 700 and 650 ◦ C was similar to the ones described above. The microstructure consisted of equiaxed primary ␣ and transformed ␤ in all the cases. Growth of primary ␣ was observed on slow cooling from 650 ◦ C. Martensite was not formed even after water quenching from 650 ◦ C due to reasons already discussed. The sample solution treated at 700 ◦ C for 1 h followed by water quenching showed globular ␣ and martensite/elongated ␣. The aging of both the water-quenched samples showed the growth ␣ while retaining the overall morphology of the water-quenched samples (Fig. 4).

The phase constituents of the heat-treated samples were also identified from the X-ray diffraction patterns shown in Fig. 5. In all the above conditions air-cooled samples showed finer structure than furnace cooled samples. 3.2. Hardness The effect of the solution treatment temperature and cooling rate on macrohardness of the samples deformed above and below ␤ transus is presented in Fig. 6. In both the cases, the variation of the hardness with cooling rates was similar. Air-cooled samples showed higher hardness than furnace cooled and waterquenched samples. Aging treatment increased the hardness markedly. The solution treatment temperature had a relatively small effect on the hardness of air-cooled and water-quenched samples. It has been reported [24,25] that the precipitation of ␣ phase in the matrix increases the hardness of the ␤ titanium alloys. In the present study, the higher volume fraction of ␣ phase and its fine distribution in the matrix increased the hardness. In general, the samples deformed above ␤ transus exhibited higher hardness than samples deformed below it. It seems that the presence of acicular or elongated ␣ increases the hardness more than globular ␣. When globular ␣ forms, elemental partitioning takes place and ␣ phase becomes lean in ␤ stabilizing

Fig. 6. Vickers hardness of the Ti–13Zr–13Nb alloy deformed at (a) 800 ◦ C and (b) 650 ◦ C and solution treated at different temperatures for 1 h followed by FC, AC, WQ and aging of the WQ sample at 500 ◦ C for 5 h.

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Fig. 7. A typical wear curve of Ti–13Zr–13Nb alloy deformed at 650 ◦ C and solution treated at 650 ◦ C for 1 h followed by AC.

alloying elements (Nb) and hence the solution hardening effect decreases. Air cooling produced a fine microstructure and hence the hardness was higher than in case of furnace cooling and water quenching. 3.3. Wear characterization During the wear tests, the temperature at the articulating surfaces increases rapidly at first and then it attains a steady value. The wear rate was estimated from the rate of height loss

or height reduction of the pin in the steady temperature range. Thus, the measurement of the height loss of the pin is not affected by thermal expansion due to gross temperature change during the test and is entirely due to wear. Three samples were tested for each set of parameters. The wear rate was calculated from the slope of the volume loss of the pin versus sliding distance plot (Fig. 7) assuming a simple linear relationship between the two and the data are presented as specific wear rate in unit of m2 /N. The slope was estimated at 99% confidence level. In case of dry wear, no significant variation of the wear rate with heat treatment was observed (Figs. 8 and 9), except for the sample deformed and solution treated at 800 ◦ C (above the ␤ transition temperature) and aged after quenching which showed considerably high wear rate. It may be recalled that in this case a large amount of martensite was formed on quenching from 800 ◦ C which decomposed predominantly into acicular ␣ during ageing at 500 ◦ C (Section 3.1). It can thus be suggested that the presence of acicular ␣ in the matrix increases the wear rate substantially. The wear rate of the all the samples increased substantially in Hank’s solution and in bovine serum. The wear loss was more severe in bovine serum than in Hank’s solution (Figs. 10–13). However, the wear behavior of the different heat-treated samples showed almost similar trend in both the lubricating conditions with furnace cooled and aged samples showing the lowest wear rate. Titanium and its alloys are considered to have poor oxidative wear resistance when ‘tribo-chemical’ reactions occur at the contact area [5]. The detail mechanism of oxidative wear has been presented in [5]. It has been reported that corrosion

Fig. 8. Wear behavior of Ti–13Zr–13Nb alloy in dry condition. Samples were deformed at 800 ◦ C and solution treated at different temperatures: (a) wear rate, (b) coefficient of friction and (c) temperature.

Fig. 9. Wear behavior of Ti–13Zr–13Nb alloy in dry condition. Samples deformed at 650 ◦ C and solution treated at different temperatures: (a) wear rate, (b) coefficient of friction and (c) temperature.

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Fig. 10. Wear behavior of Ti–13Zr–13Nb alloy in Hank’s solution. Samples were deformed at 800 ◦ C and solution treated at different temperatures: (a) wear rate and (b) coefficient of friction.

Fig. 11. Wear behavior of Ti–13Zr–13Nb alloy in Hank’s solution. Deformation temperature, 650 ◦ C. Solution treated at different temperatures: (a) wear rate and (b) coefficient of friction.

is an accelerating factor on the wear behavior [10]. Apart from its lubricating action, Hank’s solution may also cause electrochemical reaction on the surface of the titanium alloy which can accelerate wear rate. However, the influence of corrosion on wear varies with materials. The corrosion-wear as opposed to non-corrosive wear is more detrimental to the mixed phase ␣–␤ alloys than for other alloys [10]. Of all the cases studied, the

highest rate was observed when the samples deformed at 650 ◦ C and quenched from solution treatment temperatures of 700 and 650 ◦ C were tested in bovine serum. In general, faster cooling after solution treatment (i.e. air cooling and water quenching) resulted in higher wear rate in wet condition. This is particularly true for samples deformed at lower temperature (650 ◦ C) and tested in bovine serum. This can be attributed to finer microstruc-

Fig. 12. Wear behavior of Ti–13Zr–13Nb alloy in bovine serum. Samples were deformed at 800 ◦ C and solution treated at different temperatures: (a) wear rate and (b) coefficient of friction.

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Fig. 13. Wear behavior of Ti–13Zr–13Nb alloy in bovine serum. Deformation temperature, 650 ◦ C. Solution treated at different temperatures: (a) wear rate and (b) coefficient of friction.

ture and to the presence of elongated ␣. Moreover, it is possible that effect of deformation, particularly the low temperature deformation, accelerates ‘corrosion-wear’. The effect of deformation is at least partially eliminated during furnace cooling and aging and therefore the rate of ‘corrosion-wear’ is reduced. These results possibly indicate the effect of corrosion on wear and warrants an in-depth study, particularly of the sample surface, for a comprehensive understanding of the phenomenon. In general, there is no clear relationship between the hardness and wear resistance of the heat-treated samples, though the samples deformed and solution treated at 800 ◦ C show a trend in dry condition (Fig. 14). Thus, the variation in wear properties cannot be explained by hardness variation. The process of wear usually involves several mechanisms of wear which operate simultaneously [26]. Manivasagam et al. [10] have reported that adhesive wear mechanism is dominant in Ti–Zr–Nb alloy. Similarly, Long and Rack [27] have reported that adhesive wear is the principal mechanism of dry wear of Ti–6Al–4V ELI and Ti–6Al–7Nb ELI alloys articulated against hardened steel (HRc ≈ 60). Corrosion-wear tests against zircon ball in Ringer’s solution have shown adhesive wear grooves in biomedical ␤ titanium alloys [3]. In the present study, SEM examination of the wear surface in dry condition showed smeared wear surface with ridged wear scars parallel to the sliding direction (grooving wear) and superficial plastic deformation along the direction

of the sliding (Figs. 15 and 16). The dominating wear mechanism was microcutting. The abrasive scratches with transfer layer are in agreement with the observation of Choubey et al. [20]. As soon as the surface of the titanium alloy is exposed to environment it rapidly forms an oxide layer on the surface due to its high reactivity. After studying the dry sliding wear of Ti–6Al–4V alloy at a sliding speed of 0.3 m/s, Molinari et al. [28] concluded that the wear debris comprised of TiO. Li et al. [11] have reported that the main constituents of the passive film on Ti–39Nb–13Ta–4.6Zr alloy immersed in NaCl solution are Ti2 O3 and Nb2 O5 . In the present work, the presence of surface oxide could not be detected by XRD analysis because it was ultra thin, but it is expected that titanium oxides (mainly TiO2 and TiO) form at the contact surface of the materials during wear [5]. Moreover it was not possible to analyze the wear debris by XRD because the amount was very small. The repassivation capacity of the surface layer during the wear process also plays a major role during the wear process. In Hank’s solution, the wear tracks were relatively smooth with much shallower wear grooves and traces of local material removal were seen on the worn surface (Fig. 16). Rapid formation of a passive layer on the contact surface of the material and its subsequent removal during the next sliding contact may be the reason for the formation of smooth surfaces. In case of bovine serum, SEM investigation of the wear surface shows parallel deep scratches with some

Fig. 14. Hardness vs. wear rate (in dry condition) of the heat-treated Ti–13Zr–13Nb alloy deformed at (a) 800 ◦ C and (b) 650 ◦ C.

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Fig. 15. SEM micrograph of the wear track of Ti–13Zr–13Nb alloys in dry condition (a) deformed at 800 ◦ C and solution treated at 800 ◦ C for 1 h followed by AC and (b) deformed at 650 ◦ C and solution treated at 650 ◦ C for 1 h followed by WQ and aging.

Fig. 16. SEM micrograph of the wear track of Ti–13Zr–13Nb alloys in Hank’s solution (a) deformed at 800 ◦ C and solution treated at 800 ◦ C for 1 h followed by FC and (b) deformed at 650 ◦ C and solution treated at 650 ◦ C for 1 h followed by WQ.

patches of transfer layer (Fig. 17). These features suggest that the wear mechanism is mainly abrasive. The morphological analysis of the wear debris using SEM showed irregular-shaped particles with very sharp edges (Fig. 18). The fragmented particles are relatively large in Hank’s solution than in dry condition (Fig. 18a and b). The size of the debris increased further in bovine serum, indicating severe wear of the material in this medium (Fig. 18c). The concentration of protein in the lubricant has a large effect on the friction and wear of the tribological pairs used in the prosthetic joints [14]. Also, the wettability of the prosthetic

materials affects the wear behavior, though there is no definite relationship between wettability and tribological behavior of the material [14]. It has been reported that in case of metal–polymer (TiN-coated stainless steel-UHMWPE) tribological pair, the friction and wear rate is smaller in a mixture of bovine serum and Hank’s solution than in Hank’s solution [17]. The results were explained on the basis of the presence of albumin in serum which resulted in the formation of an adsorbed protein layer on the surfaces. This layer protects the surfaces of the tribological pair resulting in reduction of the interaction between the solid

Fig. 17. SEM micrograph of the wear track of Ti–13Zr–13Nb alloys in bovine serum (a) deformed at 800 ◦ C and solution treated at 650 ◦ C for 1 h followed by WQ and (b) deformed at 650 ◦ C and solution treated at 650 ◦ C for 1 h followed by WQ and aging.

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Fig. 18. SEM micrograph of the wear debris of the Ti–13Zr–13Nb alloys (a) in dry condition (650 ◦ C GR and 700 ◦ C/1 h FC), (b) in Hank’s solution (650 ◦ C GR and 700 ◦ C/1 h FC) and (c) in bovine serum (800 ◦ C GR and 800 ◦ C/1 h AC) (please note that the magnification in figure (c) is four times smaller than in (a) and (b); this has been done to show a number of particles).

surfaces, and, as a consequence, friction and wear are smaller [14,17]. Moreover, the wear mechanism was found to change from abrasive in Hank’s solution to adhesive in the mixture of bovine serum and Hank’s solution. Amaral et al. [13], on the other hand, found that both friction and wear are higher in bovine serum than in Hank’s solution in case of Si3 N4 –bioglass composite/nanocrystalline diamond (NCD)-coated Si3 N4 tribological pair. This was again attributed to the effect of protein attachment to the bearing surfaces, though this resulted in mixed lubrication regime in this case. The surfaces were characterized by regions with localized contacts between the opposing surfaces interleaved with zones where the fluid prevented direct contact. Therefore, larger contact stresses in bovine serum than in Hank’s solution tests led to increased wear. The deficiency in full fluid lubrication in the entire contact region increased the friction coefficient. The wear mechanism was abrasive in both the lubricants. These apparently contradictory results together with some more similar results presented in [29] suggest that wear behavior in bovine serum depends on the material characteristics of the tribological pair with solid/lubricant interfacial energy playing a major role. Higher wear and friction in bovine serum than in Hank’s solution in the present investigation is consistent with the results of Amaral et al. [13] and is possibly due to incomplete fluid lubrication. Primarily abrasive wear mechanism was operative (Figs. 16 and 17) in both the solutions, but it was more severe in bovine serum. Fig. 19 shows the microstructure of the cross-section of the worn specimen where no plastically strained subsurface is

observed. The applied normal load was low (50 N) and this was not enough for any subsurface deformation. The hardness of the specimen surface did not change after wear in Hank’s solution. After dry wear testing, hardness of the wear surface could not be measured because the impressions were not distinguishable. In Hank’s solution, discrete patches of corrosive layer were observed on the wear surface. In all the cases, the coefficient of friction (measured in the steady temperature region) varied from 0.25 to 0.46. The temperature varied from 82 to 112 ◦ C in dry condition and from 42 to 45 ◦ C in both the media. The cooling effect of the lubricants is clearly evident which prevented any significant temperature rise during wet wear tests. Relatively high coefficient of friction (0.4–0.75) was observed for titanium sliding against Al2 O3 at room temperature [5]. Mercer and Hutchings [30] have reported a coefficient of friction of about 0.6 during abrasive wear of Ti–6Al–4V against the silicon carbide and aluminium oxide abrasive papers. Choubey et al. [20] have reported that the coefficient of friction for Ti–13Zr–13Nb alloy against SAE 52100 bearing steel (high-carbon chromium alloy steel) in Hank’s solution is 0.46, which is consistent with the present work. In bovine serum, coefficient of friction varied from 0.37 to 0.5, which is higher than that in Hank’s solution. This, as already discussed, may be due to lack of full fluid lubrication in the entire contact region as a result of the protein attachment to the contact surfaces [17]. The presence of corrosive layer on the wear surface (Fig. 19) suggests that the corrosion accelerated the wear of the samples in Hank’s solution.

Fig. 19. Microstructure of the subsurface region of the worn samples of the Ti–13Zr–13Nb alloys deformed at 650 ◦ C and solution treated at 700 ◦ C/1 h WQ followed by aging at 500 ◦ C for 5 h (a) in dry condition and (b) in Hank’s solution.

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4. Summary and conclusions I. Ti–13Zr–13Nb alloy was hot worked and heat treated under different conditions to produce a variety of microstructures. II. The microstructure of the heat-treated alloy consisted of ␣, ␤ and martensite. Hot working below ␤ transus and subsequent heat treatment at different temperatures produced primary ␣ of globular morphology. In almost all the cases, aging of the water-quenched samples resulted in globular ␣. However, the sample deformed at 800 ◦ C and solution treated at 800 ◦ C for 1 h followed by water quenching showed partially acicular ␣ after aging. III. In all cases, the variation of the hardness with cooling rates was similar. Air-cooled samples showed higher hardness than furnace cooled and water-quenched samples. Aging treatment increased the hardness markedly. Precipitation of ␣ phase and its fine distribution in the matrix increased the hardness. IV. In dry condition, there was no significant variation of the wear rate of the heat-treated samples except for the sample deformed above ␤ transus temperature and solution treated at 800 ◦ C followed by water quenching and aging. Higher wear rate in this case is possibly because of the presence of acicular ␣. V. In simulated body fluids, wear rate of the all samples increased substantially. The wear rate was higher in bovine serum than in Hank’s solution. In bovine serum, the effects of proteins attachment to the bearing surfaces and in Hank’s solution a combined effect of abrasion and corrosion seem to have increased the wear rate of the material. VI. The wear process was characterized by ridged wear scars parallel to the sliding direction with superficial plastic deformation along the direction of the sliding and by smeared wear surface. The wear mechanism is primarily abrasive. Acknowledgement The authors would like to acknowledge Nuclear Fuel Complex, Hyderabad, for supplying the raw materials for this research work. References [1] Y.L. Hao, M. Niinomi, D. Kuroda, F. Fukunaga, Y.L. Zhou, R. Yang, A. Suzuki, Young modulus and mechanical properties of Ti–29Nb–13Ta– 4.6Zr in relation to ␣ martensite, Metall. Mater. Trans. A 33 (2002) 3137–3144. [2] H. G¨ulery¨uz, H. Cimeno˘glu, Effect of thermal oxidation on corrosion and corrosion-wear behavior of a Ti–6Al–4V alloy, Biomaterials 25 (2004) 3325–3333. [3] M. Niinomi, D. Kuroda, K.I. Fukunaga, M. Morinaga, Y. Kato, T. Yashiro, A. Suzuki, Corrosion wear fracture of new ␤ type biomedical titanium alloys, Mater. Sci. Eng. A 263 (1999) 193–199. [4] M. Geetha, U.K. Mudali, A.K. Gogla, R. Asokamani, B.D. Raj, Influence of microstructure and alloying elements on corrosion behavior of Ti–13Nb–13Zr alloy, Corros. Sci. 46 (2004) 877–892. [5] M. Long, H.J. Rack, Titanium alloys in total joint replacement—a materials science perspective, Biomaterials 19 (1998) 1621–1629.

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