Surface mechanical attrition treatment of low modulus Ti-Nb-Ta-O alloy for orthopedic applications

Materials Science & Engineering C 110 (2020) 110729

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Surface mechanical attrition treatment of low modulus Ti-Nb-Ta-O alloy for orthopedic applications

T

Srijan Acharyaa, Arpana Gopi Panickera, Vasanth Gopalb, Shaurya Singh Dabasa, ⁎ Geetha Manivasagamb, Satyam Suwasa, Kaushik Chatterjeea, a b

Department of Materials Engineering, Indian Institute of Science, Bangalore 560012, India Centre for Biomaterials, Cellular and Molecular Theranostics (CBCMT), Vellore Institute of Technology, Vellore, TN 632014, India

A R T I C LE I N FO

A B S T R A C T

Keywords: Orthopedic biomaterials Ti-Nb-Ta-O β titanium alloys Surface engineering Corrosion Fretting wear

Surface mechanical attrition treatment (SMAT) is recognized as a surface severe plastic deformation (SPD) method that is effective in improving the surface-dependent mechanical and functional properties of conventional metallic biomaterials. In this study, we aimed to systemically investigate the effect of SMAT on the physical, electrochemical, tribological and biological performances of a newly developed low modulus β Ti-NbTa-O alloy with two different microstructures, namely, single phase β-treated and dual phase β + α aged. The microhardness results showed considerable hardening for the β-treated condition due to formation of deformation substructures; that was associated with increased corrosion resistance resulting from a stronger and denser passive layer on the surface, as revealed by Tafel polarization, impedance studies and Mott-Scottky plots. The wear volume loss during fretting in serum solution was found to decrease by 46% while friction coefficient decreased only marginally, due to presence of a harder and more brittle surface. In the β + α condition of the alloy, minimal hardening was observed due to coarsening of the precipitates during SMAT. However, this also reduced the number of α-β interfaces, which in turn minimized the tendency for galvanic corrosion resulting in lower corrosion rate after SMAT. Wear resistance was enhanced after SMAT, with 32% decrease in wear volume loss and 21% decrease in friction coefficient resulted due to improved ductility on the surface. The attachment and growth of osteoblasts on the alloys in vitro were not affected by SMAT and was comparable to that on commercially pure Ti. Taken together, these results provide new insights into the effects of surface SPD of low modulus β- Ti alloys for orthopedic applications and underscore the importance of the initial microstructure in determining the performance of the alloy.

1. Introduction One of the key limitations associated with the current generation of high strength implant materials for load bearing applications such as Ti6Al-4V, 316 stainless steels and CoeCr alloys, is their high elastic modulus (> 100 GPa). The mismatch in modulus between the implant and human bone (< 30 GPa) induces stress-shielding effect, which in turn leads to osteopenia around the implant and can result in failure of the implant due to poor osseointegration and eventual loosening [1]. In this context, titanium alloys primarily consisting of metastable β-phase have shown great promise, as this phase has relatively lower elastic modulus (< 80 GPa) than that of the α-phase (> 100 GPa) [1], which is thermodynamically stable below 882 °C in pure titanium. These alloys need to have sufficient content of β-stabilizing elements to stabilize the β-phase. However, the elastic modulus is also a function of the type and

⁎

amount of β-stabilizer present in the material. Previous studies have reported that the lowest modulus for the β-phase is obtained for composition with an average electron per atom ratio (e/a) of 4.2 to 4.24 [2]. Nb and Ta impart the lowest modulus values for this composition range [2–3]. These two elements have also been found to be non-toxic for human body, unlike other elements such as Al, V, Ni, which are present in conventional implants [4]. Thus, we designed a low modulus Ti-34Nb-2Ta-0.5O (wt%) alloy [5–6]. This alloy is similar in composition to another Ti-35.7Nb-2Ta-3Zr-0.3O (wt%) alloy, popularly known as gum metal, that is also been investigated for orthopedics [7–8]. The absence of Zr in the Ti-34Nb-2Ta-0.5O (wt%) alloy, in comparison to the gum metal, does not affect the stability of the β-phase with respect to formation of high-modulus ω-phase, due to additional oxygen (0.5 wt %) being present [5]. In addition, as an interstitial alloying element, oxygen was also found to strengthen the alloy.

Corresponding author. E-mail address: [email protected] (K. Chatterjee).

https://doi.org/10.1016/j.msec.2020.110729 Received 5 August 2019; Received in revised form 19 December 2019; Accepted 3 February 2020 Available online 04 February 2020 0928-4931/ © 2020 Elsevier B.V. All rights reserved.

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In our previous study, we reported a deterioration in the fretting wear behavior after age-hardening of the Ti-Nb-Ta-O alloy despite the increase in tensile strength which was ascribed to increased tribocorrosion and formation of a harder transfer layer [24]. The single phase β microstructure of the alloy exhibited better wear properties but suffered from lower tensile strength, which is critical for designing load bearing implants. Toward enhancing the performance of the alloy for orthopedic applications, the effects of surface SPD by SMAT on the mechanical and functional responses were investigated on both single phase β-annealed and dual phase aged conditions of the Ti-Nb-Ta-O alloy. Whereas electrochemical behavior was characterized by polarization and impedance studies, the fretting wear response were recorded in terms of friction studies against ZrO2 balls in protein containing medium that mimic in vivo conditions. The effect on cytocompatibility, if any, after SMAT was measured by culturing osteoblast. The variation in these properties have been systematically investigated by analysis of the microstructures of the SMAT processed samples, using scanning electron microscopy (SEM) and electron backscattered diffraction (EBSD) studies.

The surface properties of a biomaterial play a critical role in determining the performances of an orthopedic implant. These properties determine its corrosion behavior and wear resistance thereby modulating release of metal ions, which cause toxicity, inflammation and bone resorption, eventually leading to implant loosening and failure [1,9–10]. Furthermore, a hardened surface layer of the implant not only improves the wear resistance, but can also delay the initiation of fatigue failure [11–12]. Finally, the biological response of the implant is also affected by its surface properties such as roughness, wettability, surface charge and composition [10,13]. Titanium alloys are known for their poor wear resistance [14–15]. Moreover, osseointegration of titanium alloy surfaces is limited by their “bio-inert” surface characteristics [16]. The most common surface modification strategies for orthopedic titanium alloys have been based on chemically coating with hard materials (e.g. titanium nitride, diamond-like carbon) or calcium phosphate-based bioactive materials (hydroxyapatite, calcium phosphate) [17–19]. However, these methods have some inherent disadvantages such as relatively poor bond strength between the coating and substrate due to difference in their chemical composition, and higher surface roughness produced by these processes which is detrimental for fatigue resistance. Apart from these processes, surface-based severe plastic deformation (S2PD) methods have also been attempted for these materials to obtain harder and bioactive surfaces without compromising the chemical composition and surface roughness of the material. Among the different S2PD techniques, the surface mechanical attrition treatment (SMAT) has been the most popular one. In SMAT, typically nanocrystalline grains are generated owing to severe plastic deformation (SPD) of the surface [12,20]. Consequently, SMAT was observed to impart several beneficial effects such as improved surface hardness and fatigue strength as well as superior cellular response to commercially pure Ti (cp-Ti) [12] and stainless steel [11]. A few studies have been reported on the effect of SMAT on surface microstructure and properties in β-Ti alloys. Nanocrystallized grains were obtained in a β-solution treated Ti–30Nb–8Zr–0.8Fe (wt%) alloy by SMAT [21], whereas Han et al. reported that SMAT of aged βTi–25Nb–3Mo–3Zr–2Sn alloy results in a reverse α → β transformation that finally leads to the formation of nanocrystallized β grains [22]. Enhancements in corrosion resistance as well as in response to mesenchymal stem cells and osteoblasts have been reported after SMATprocessing of β-Ti–25Nb–3Mo–3Zr–2Sn and β-Ti–32Nb–2Sn alloys [13,23]. These reports indicate that SMAT can be effective in improving the performances of biocompatible β-Ti alloys. However, given the wide variety of β-Ti alloys, there is need for continued investigation and SMAT of Ti-Nb-Ta-O based alloys has not been reported earlier. Importantly, there are no reports on the effect of SMAT on wear properties for any β-Ti alloy, which is essential considering their use in load bearing orthopedic and dental implants where fretting is widely noted.

2. Materials and methods 2.1. Material and processing Cast pancake of Ti-34Nb-2Ta-0.5O alloy with ≈96 mm diameter and ≈12 mm thickness was produced by vacuum arc melting using a Tungsten electrode. Radial strips of 50 mm length and 20 mm width were extracted from the pancake and subjected to homogenization above the β-transus temperature at 900 °C for 1 h and then hot rolled at the same temperature to a thickness of ≈5 mm (strain εp = 0.9) at a true strain per pass of 0.2 (εp = 0.2) to break the as-cast microstructure [5,24]. Subsequently, hot rolled plates were β-solutionized at 900 °C for 0.5 h and then quenched in water. These samples were cold rolled up to a 70% thickness reduction and were further annealed at 900 °C for 0.5 h to obtain recrystallized microstructure consisting of equiaxed β-grains. Before hot rolling and annealing, samples were coated with Delta glaze to prevent oxidation in the furnace. Some of these annealed samples were aged at 500 °C for 4 h to obtain precipitation of α in the β-matrix. Normal surfaces of the annealed samples and the aged samples, hereafter to be referred to as STQ and Aged, respectively, were mechanically polished up to P3000 paper followed by electropolishing at 40 V for 20 s using A3 solution in a Struers Lectropol-5 machine, prior to SMAT processing. SMAT was carried out at a frequency of 25 Hz for 1 h using a custom-designed unit (developed by Cosmic Industrial Laboratories Limited, Bengaluru, India) [Fig. 1]. 500 numbers of hardened steel balls, each of 4.75 mm diameter, were used for this study.

Fig. 1. Set-up for the surface mechanical attrition treatment (SMAT) along with schematic showing different parts. 2

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The working distance between base of the chamber and sample surface was set at 20 mm. The SMAT-processed STQ and Aged samples are hereafter referred to as STQ-SMAT and Aged-SMAT.

tests were done at room temperature (≈25 °C).

2.2. X-ray diffraction and microstructure

The friction and wear performance of SMAT-processed surfaces was assessed by ball-on-plate fretting wear measurement in a fretting wear machine (Ducom, India), using ZrO2 balls of 6 mm diameter against the sample surfaces, immersed in a 20% fetal bovine serum (FBS) solution (pH = 7.4) [24]. The normal load and relative displacement stroke, between the ball and sample, were fixed at 5 N and 200 μm, respectively, and the temperature was maintained at 37.4 °C. The tests were performed at a frequency of 10 Hz till 50,000 cycles. After the test, the worn surfaces were characterized using SEM.

2.5. Tribological investigation

The X-ray diffraction patterns were recorded on the normal planes of samples with 7 × 7 mm2 area and 1 mm thickness, before and after SMAT. A PanAnalytical X'pert Pro X-ray diffractometer with Cu-Kα radiation was employed for recording the patterns, at a scan speed of 0.011° s−1 and step size of 0.03°, with the 2θ range of 30°–90°. Transverse cross-sections of these samples were polished and then etched by Kroll's reagent (2% HF, 6% HNO3, 92% H2O) for 30 s, for primary microstructural examination performed using an optical microscope (Zeiss) as well as a scanning electron microscope (ESEM, Quanta FEI). The surface microstructures were further analyzed using electron back-scattered diffraction (EBSD) analysis of the cross-sections of samples polished with colloidal silica.

2.6. Surface characterization Static water contact angle on the alloy surfaces were measured on at least three replicates for each processing condition using a contact angle goniometer (OCA 15EC, Dataphysics) by placing 1 μl of ultrapure water (Sartorius Arium) droplet on the surface. Effect of SMAT on surface roughness of STQ and Aged samples was also estimated using non-contact optical profilometer (Talysurf CCI, Hobson) for a surface area of 1660 μm × 1660 μm at a magnification of 10×.

2.3. Mechanical tests Elastic moduli of the STQ, Aged, STQ-SMAT and Aged-SMAT samples were measured from the indentation-based technique using an instrumented indentation testing machine (CSM). In this method, the slope of the unloading curve obtained from the load-indentation plot is used to estimate the modulus of the material using the equation suggested by Oliver and Pharr [25,26]:

2.7. Cell attachment and proliferation Ti-Nb-Ta-O alloy and cp-Ti sample surfaces of 7 × 7 mm2 area were prepared, as described above, for cell studies. Samples were then sterilized in 70% ethanol and under UV for 30 min and placed individually in wells of a 48 well tissue culture plates. 300 μL of cell suspension containing 5 × 103 MC3T3-E1 subclone 4 mouse calvarial pre-osteoblasts (ATCC) was added to each well. The medium consisted of αminimum essential medium (α-MEM) supplemented with 10% (v/v) FBS (Gibco, Life Technologies) and 1% (v/v) Penicillin-streptomycin (Sigma-Aldrich). Cells were passaged using trypsin-EDTA and up to 31th passage were used. Quantitative assessment of cell viability was done at 1 d, 4 d and 7 d post-seeding, using the WST assay (Invitrogen). Cells were stained and imaged by fluorescence microscopy to study their morphology. For each time interval, three replicates (n = 3) of each sample were used for the assay while two were used for imaging. WST assay was performed after replacing the medium with the working solution made of 10 μl WST-1 reagent in 100 μl of culture medium, followed by incubation in 5% CO2 at 37 °C for 3 h. A plate reader (Biotek) was used to measure the absorbance at 440 nm. For statistical analysis of the WST assay results, analysis of variance (ANOVA) with Tukey's test was performed and differences were considered significant at p > 0.05. For imaging, cells were fixed with 3.7% formaldehyde for 30 min (at 25 °C) and permeabilized with 0.2% Triton X (Sigma Aldrich) for 5 min. Thereafter, they were stained with 25 μg/mL Phalloidin Alexa Fluor 546 (Invitrogen) for 15 min (for actin filaments) and 0.2 μg/mL DAPI (Invitrogen) for 5 min (for nuclei) at 25 °C. Stained cells were imaged with an inverted fluorescence microscope (Olympus).

(1 − νi2 ) 1 (1 − ν 2) = + Eeff E Ei where Eeff is the effective Young's modulus obtained from the slope of unloading curve, E and Ei are Young's moduli of specimen and indenter, respectively, and ν and νi are Poisson's ratios for specimen and indenter, respectively. The Surface hardness values of the annealed and SMAT-processed samples were measured by Vickers micro-indenter at the indentation load and dwell time of 25 gf and 10 s, respectively. Microhardness profiles were also measured on the polished cross-section of SMATprocessed samples to monitor the through-thickness variation in hardness from the SMAT-surface. 2.4. Electrochemical studies The effect of surface treatment on electrochemical behavior of the single-phase and dual-phase β‑titanium alloys was examined by Tafel extrapolation and electrochemical impedance study (EIS) techniques in simulated body fluid (SBF) solution, using a standard three electrode potentiostat (CHI604E, C.H. Instruments) with Pt as counter electrode and saturated calomel electrode (SCE) as reference 5. The composition of the SBF solution is: 8.0 g/L NaCl, 0.35 g/L NaHCO3, 0.224 g/L KCl, 0.228 g/L K2HPO4.3H2O, 0.30 g/L MgCl2.6H2O, 40 mL/L 1 M HCl, 0.278 g/L CaCl2, 0.071 g/L Na2SO4, 6.0 g/L (CH2OH)3CHN2 [5,27], with ion concentration and pH (=7.4) levels similar to that of human blood plasma. The SMAT-processed samples were prepared by the same metallographic techniques, as described above, and then were ultrasonically cleaned. The size of each sample was kept at 7 × 7 × 1 mm3. The samples were first immersed in the in SBF for 3 h to stabilize the open circuit potential (OCP) prior to Tafel extrapolation and EIS measurements. The parameters for Tafel plots were kept at scan rate of 2 × 10−4 V s−1 and voltage range of −0.6 V to +0.4 V. The EIS spectra were measured at OCP using a sinusoidal signal with 5 mV amplitude in the frequency range 105–10−1 Hz. The Mott-Schottky plots, after 3 h stabilization in OCP, were generated within a voltage range 1.0 V to −1.0 V at a step size of 50 mV using a sinusoidal signal with frequency of 1 kHz and 0.005 V amplitude. All the electrochemical

3. Results and discussion 3.1. Microstructural features of the starting material The X-ray diffraction (XRD) pattern obtained from STQ contains βpeaks only, whereas that from aged sample shows peaks from α as well (Fig. 2). A representative optical microstructure of STQ, which consists of equiaxed recrystallized grains of β-phase is shown in Fig. 3a. Fig. 3b shows a scanning electron micrograph of aged sample wherein uniform precipitation of α-platelets is seen along the grain boundary (intergranular α) as well as inside the β-grains (intragranular α) formed in the β-matrix. 3

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Fig. 2. X-ray diffraction patterns of STQ (bottom) and Aged samples (top). Aged condition shows peaks for both α- and β- phases whereas only β-peaks are seen in STQ.

Fig. 4. XRD patterns of STQ-SMAT (bottom) and Aged-SMAT samples (top).

arrows) running from the surface toward the interior are clearly visible. These are identified as microbands formed by plastic deformation during SMAT. The Inverse Pole Figure (IPF) map obtained from the EBSD scans is shown in Fig. 6a. The color variation in the IPF map indicates the changes in local misorientation caused by plastic deformation. Fig. 6b is the plot of misorientation angle as a function of distance from the surface both point-to-point and point-to-origin. Although the point-to-point misorientation is small, it increases when point-to-origin misorientation is considered. This can arise due to gradual misorientation changes from surface to bulk. Apart from this, there are also large number of misindexed spots (black), that is, spots with a confidence index (CI) value < 0.1. These spots are mostly arranged in a linear manner, which indicates that these correspond to the microbands formed by strain localization. The large strain localization along these bands results in a heavily distorted crystal structure inside them leading to their lower CI values. The evolution of microbands is a common

3.2. Microstructural evolution after surface treatment Fig. 4 shows the XRD patterns of the SMAT-processed samples (hereafter referred to as STQ-SMAT and Aged-SMAT). STQ-SMAT shows peaks only from the β-phase, but the relative intensities of the peaks are different when compared to STQ (Fig. 2). The peaks also show broadening after SMAT, which may be ascribed to plastic deformationinduced reduction in crystallite size and/or increased micro-strain. The XRD pattern for Aged-SMAT shows peaks from both β and α phases. The relative intensities in this case are also different from that of the corresponding aged samples (Fig. 2). The cross-sectional microstructure of STQ-SMAT is shown in Fig. 5. The region affected by SMAT is marked by the dashed line near the top edge. Some etch pits are clearly observed in the sample. Although grains are not clearly revealed here, several bands (marked by blue

Fig. 3. (a) Optical micrograph of STQ samples showing equiaxed β-grains, and (b) scanning electron micrograph showing α-precipitates (seen in the oval marks) in βmatrix. 4

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Fig. 5. Cross-sectional SEM micrographs of STQ-SMAT; black dashed line indicates the transition from deformed layer to the bulk and the blue arrows show the direction of bands. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

Fig. 6. (a) IPF map obtained from EBSD analysis of transverse cross-sectional plane near the SMAT-processed surface (bottom) of STQ-SMAT sample; (b) misorientation angle plot as a function of distance from the surface showing gradual misorientation build up from surface.

Fig. 7. (a) Low magnification and (b) high magnification SEM micrographs of Aged-SMAT; the microbands are marked by arrow whereas α-platelets are encircled.

feature in deformation of β‑titanium alloys and several other materials [28–30]. The mechanism of formation of microbands has been well investigated and has been primarily attributed to splitting of dense dislocation walls, which separate the cell blocks within the grains [11,29,31]. The dislocation boundaries along with the microbands

results in fragmented structure of grains. However, there is no clear evidence of nanocrystallization after SMAT. In literature, nanocrystallized grains have been reported on surface of several materials as a result of SMAT [12,20,32–33]. In case of the β-Ti alloys, the formation of nanocrystallized grains by severe plastic deformation has been 5

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attributed to be dependent upon the stability of β-phase [34]. For alloys with lower stability, the formation of nano-scale orthorhombic martensite (α″) during deformation assists in refinement of grains by causing barriers to dislocation slip. In comparison, for alloys with higher β-stability, martensite is absent in the microstructure and deformation takes places by conventional stages of dislocation substructure and cell formation which is not followed by formation of nanocrystals. In the present case, the presence of oxygen increases the stability of the β-phase to minimize the possibility of martensitic transformation. This higher stability can be a possible reason for the absence of nanocrystallized grains. In addition, the energy of the impact also affects the microstructural evolution depending upon the hardness of the material. The microstructure of Aged-SMAT is shown in Fig. 7. In Fig. 7a, α platelets are seen near the surface along with the micro-bands. At higher magnification, microbands are seen to be aligned along the arrow marked in the figure (Fig. 7b). The morphology of these platelets is similar to those observed in the aged samples (Fig. 4). However, the platelets are larger after SMAT. This indicates that there is coarsening of the precipitates due to SMAT. Although intriguing and unexpected, similar observations were noted in SMAT-processed β Ti-32Nb-2Sn alloy [13]. Coarsening of precipitate is a diffusion-controlled process, which is facilitated at higher temperature. However, there have also been reports of dynamic precipitate coarsening during deformation at ambient temperature [35–36]. This again is ascribed to deformation-induced dissolution of precipitates prone to shear, which can cause the surrounding matrix to become supersaturated with solute atoms. This can lead to growth of the precipitate particles. In fact, the metastable β Ti–25Nb–3Mo–3Zr–2Sn alloy has been found to show α → β phase transition during SMAT in the aged condition [22]. In the present study, no evidence of SMAT-induced dissolution of precipitates has been detected from the XRD or scanning electron microscopy. However, the coarsening of precipitates indicates that a similar mechanism is likely to operate here.

similar to that of the bulk hardness indicating that the effect of SMAT is essentially limited to ≈25 μm from the surface. In general, the hardening imparted by SMAT has been attributed to nanocrystallization of grains on the surface by surface SPD [33,37], although substantial hardening results from dislocation hardening as well [20]. In this case, the EBSD reveals fragmentation of grains by the microbands and presence of dislocation substructures (Fig. 6). The presence of dislocation boundaries in the grains on the surface results in a cell structure giving rise to work hardening. The microbands present in these regions further contribute to strengthening by serving as barriers to dislocation motion. With increase in the distance from the surface, the defects decrease in number and thus hardness reduces. The present results indicate that SMAT can also impart considerable hardening of the surface of β-Ti alloys even without producing nanostructured grains. Surface hardness for aged samples remains unchanged after SMAT. The micro-hardness profile also shows negligible change of hardness from surface to bulk. This is in sharp contrast to the results for STQ. The presence of microbands in Aged-SMAT samples indicates the occurrence of plastic deformation. However, the presence of deformationinduced defect structure does not account for the overall strengthening. This can be ascribed to the reduced number of the α-β interfaces in the SMAT-processed region thereby reducing the precipitation hardening effect. Moreover, any dissolution of α precipitates could further reduce the volume fraction of α phase. In this alloy, the presence of oxygen leads to strengthening of the α particles by enriching interstitial oxygen, which is an α-stabilizer [38]. This makes deformation of α phase more difficult and hence, only β-phase can be deformed. The strengthening of the β-matrix compensates for this by coarsening and possible dissolution of the α phase. 3.4. Effect of SMAT on surface wettability The surface wettability for the different samples, measured by static water contact angle are presented in Table 1. These values are very close with minimal differences found before and after SMAT processing for both the STQ and Aged samples. Surface wettability of a material is a general indication of the energy of a surface, which increases with decreasing water contact angle. The presence of higher number of grain boundaries and higher dislocation densities in nanocrystallized surface formed by SMAT can increase the energy, thus resulting in a lower water contact angle. However, in this case, the minimal differences between the water contact angles indicates no variation in surface energy caused by SMAT. As surface energy is affected by surface roughness as well, Table 1 also lists the average surface roughness values which were found to increase marginally after SMAT. This effect was found in both STQ and aged conditions. It is reported in literature that SMAT can increase the surface roughness [13]. It is to be noted that the alteration in surface roughness by SMAT depends on the surface conditions prior to SMAT. In the present study, samples were made

3.3. Effect of surface treatment on elastic modulus and hardness The elastic modulus values for (mean ± S.D. for n = 3) for STQ, STQ-SMAT, Aged and Aged-SMAT has been found to be 64 ± 4 GPa, 66 ± 3 GPa, 86 ± 4 GPa, 83 ± 5 GPa, respectively. These values, indicate that SMAT has not affected the bulk elastic modulus of the samples. The Vickers' micro-hardness value on the normal surface of STQ and STQ-SMAT are 290 ± 7 Hv and 351 ± 5 Hv, respectively, whereas that for the Aged and Aged-SMAT, the values are 334 ± 5 Hv and 348 ± 6 Hv, respectively. The cross-sectional micro-hardness profiles obtained along the thickness of the SMAT processed samples are presented in Figs. 8(a and b), for STQ-SMAT and Aged-SMAT, respectively. For the hardness at depth = 0 μm (surface), surface hardness values have been used. The hardness decreases with increasing distance from the surface. Beyond a depth of ≈25 μm, the hardness is

Fig. 8. Microhardness profiles as a function of distance from surface to depth in (a) STQ-SMAT and (b) Aged-SMAT. 6

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two layers: a porous outer layer followed by a barrier layer inside. These two layers consist of resistances (RP and RB for the porous and barrier layers, respectively) in parallel to constant phase elements (CPE: QP and QB). RS is the resistance corresponding to the electrolyte solution. The CPE behaves like a non-ideal capacitor and it is represented by the following equation:

Table 1 Static water contact angle and surface roughness values. Condition

Contact Angle (°)⁎

Ra (nm)⁎

STQ STQ-SMAT Aged Aged-SMAT

62 65 67 69

74 95 81 91

⁎

± ± ± ±

3 3 2 2

± ± ± ±

7 16 3 12

Z = [Q (jω)n]−1

Mean ± S.D. for n ≥ 3.

(1)

where Z and Q represent the impedance and constant phase element, j is the imaginary unit and ω denotes angular frequency of the sinusoidal signal. The ‘n’ value in the exponent determines the behavior of the CPE. For n = 1, it is a capacitor and for n = 0, it is a resistance. Values for all the equivalent circuit components used to fit the impedance plots are summarized in Table 3. The Bode impedance plots are comparable in nature for the two conditions. In both cases, the impedance remains low and nearly constant in the high frequency range. This region is dominated by the solution resistance. With further reduction in the frequency, the impedance linearly increases primarily due to the capacitive response of the passive oxide layers. The phase angle plots in Fig. 10a are also similar in nature for the two conditions. Here, the phase angle increases with decreasing frequency and then remains close to 90° due to increasing contribution from the capacitance of the oxide layer. Unlike STQ-SMAT, the phase angle drops again in STQ as frequency is further lower, which is generally attributed to the resistance of the passive film. Both QP and QB become lower after SMAT. The RB have also increased after SMAT while there is a decrease in the RP value after SMAT. However, the RB values themselves are much higher as compared to RP indicating that the protection is dominated by the inner layer thereby making the overall oxide layer more protective after SMAT. Mott-Schottky plots for the two conditions are shown in Fig. 11. The positive slope for both the conditions indicates that the oxide layer is ntype in character. The charge carrier density of the oxide layer is related to the slope of the following equation:

Fig. 9. Tafel plots obtained from polarization study showing lower corrosion rates in STQ-SMAT and Aged-SMAT, compared to STQ and Aged samples, respectively.

essentially flat by electro-polishing such that subsequent SMAT processing marginally increased the roughness. It is likely that any increase in the surface energy caused by plastic deformation seems to have been negated by this minor change in the surface roughness.

1 2 ⎞⎛ kT ⎞ E − Efb − =⎛ e ⎠ Csc2 ⎝ εr ε0 eND ⎠ ⎝ ⎜

⎟

(2)

where, CSC is the capacitance of space charge region, ND is the donor density, εr is the relative permittivity of the oxide layer and ε0 is the permittivity of vacuum, E is the applied potential, Efb is the flat band potential, k is Boltzmann's constant, T is the temperature, and e is the charge of an electron. The estimation of the slopes of the linear part of the curves reveals > 50% reduction of the charge carrier density after SMAT. The lower charge carrier density can explain the improvement in the corrosion resistance observed from Tafel plot, as lower charge carrier density helps in inhibiting the dissolution of ions [12]. It has been reported that SPD imparts large number of defects in the material leading to formation of stronger passive oxide on the surface [40,41]. The lower charge carrier density accounts for the higher stability of the oxide layer. Thus, a simultaneous decrease in corrosion rate as well as critical current density for passivation is observed for the SMAT-induced SPD on the surface.

3.5. Electrochemical behavior The representative Tafel plots showing polarization behavior of different samples in SBF solution are compiled in Fig. 9. Icorr, which measures the corrosion rate, is the lowest for STQ-SMAT followed by STQ, Aged-SMAT and Aged (Table 2). The critical current for passivation (Ipass) also increases in this same order, in fact more prominently than the Icorr values. These results indicate that corrosion resistance and tendency to passivate increase for both STQ and Aged after SMAT. The higher Icorr values for both Aged and Aged-SMAT compared to STQ can be ascribed to the two-phase microstructure, which yields a galvanic couple and non-uniform oxide layer on the surface [39]. Fig. 10a shows the Bode plots obtained experimentally from the EIS studies as well as the simulated ones for STQ and STQ-SMAT. The Bode impedance is plotted logarithmically while (−phase angle) is plotted against frequency (logarithmic scale). The equivalent circuit model used for simulating the experimental data is shown in Fig. 10b. According to this model, the passive film on the alloy surface comprises of

3.6. Tribological behavior Fig. 12 shows representative plots of change in the coefficient of friction (COF) with respect to the number of cycles of fretting wear for

Table 2 Tafel plot parameters. Parameters −3

2 ⁎

μA/cm ) Icorr (×10 Ecorr (V vs SCE)⁎ ⁎

STQ

STQ-SMAT

Aged

Aged-SMAT

21.0 ± 3.0 −0.22 ± 0.02

7.4 ± 2.7 −0.14 ± 0.03

152.0 ± 9.0 0.32 ± 0.04

69.9 ± 0.3 0.26 ± 0.05

Mean ± S.D. for n ≥ 3. 7

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Fig. 10. Measured and simulated (a) Bode phase and Bode impedance plots of STQ and STQ-SMAT; (b) equivalent circuit used to fit the plots. Table 3 Values of the elements obtained from impedance spectroscopy of STQ and STQSMAT. Elements

STQ

STQ-SMAT

RS (ohm cm2) RP (ohm cm2) QP (μF cm−2) n (for QP) RB (×105 cm2) QB (μF cm−2) n (for QB)

28 1924 18 0.83 4 6 0.75

24 68 12 0.90 18 3 0.90

Fig. 12. Representative plots of coefficient of friction (COF) against number of wear cycles. Table 4 Fretting wear properties. Condition

COF⁎

STQ STQ-SMAT Aged Aged-SMAT

0.46 0.44 0.68 0.56

⁎

Wear Volume Loss (mm3)⁎ ± ± ± ±

0.04 0.06 0.03 0.05

0.76 0.52 1.13 0.85

± ± ± ±

0.09 0.11 0.12 0.10

Mean ± S.D. for n ≥ 3.

indentation further enlarges with increasing number of cycles, primarily along the sliding direction. With further continuation of the abrasive action cracks are initiated eventually leading to pull-out of worn material. This process is facilitated by electrochemical dissolution. The initial pressure leads to damage in the surface oxide film, which is brittle in nature. This exposes the metal to the solution, which works as a tribo-corrosive environment for localized corrosion to help in removal of material [42]. In general, the diameter of the wear scars is proportional to the depth of the wear scar, which primarily determines the amount of material removed during the process. The wear volume loss (V) is commonly taken as measure of wear rate. The approximate values of V for the different conditions were calculated using the following equation [43,44]:

Fig. 11. Mott-Schottky plots (Impedance vs. potential) for STQ and STQ-SMAT in SBF, showing higher positive slope for STQ-SMAT.

the four different conditions. All the curves are characterized by a runin step with sharp rise followed by a drop in the COF value and a steady-state region when the COF remains nearly constant. The average COF values after 50,000 cycles are summarized in Table 4. COF of Aged is higher than for STQ, whereas that for STQ-SMAT and Aged-SMAT are similar to that of STQ. Micrographs of the wear scars produced by the fretting process are nearly elliptical in shape with the major axes lying along the sliding direction (Fig. 13). The wear scars are characterized by the presence of wear tracks along the major axis and regions showing material pull-out. This is characteristic of gross-slip wear mechanism. The wear process initiates with a Hertzian contact pressure acting at the interface between the sample surface and the counterpart ZrO2 ball creating a depression on the surface during initial cycles by plastic deformation. This

h ⎞4 V = πRh2 ⎛1 − 4 R⎠ ⎝

(3)

where, R is radius of the counterpart (ZrO2 ball) and h is the depth of 8

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Fig. 13. Scanning electron micrographs of wear scars in (a) STQ, (b) STQ-SMAT, (c) Aged and (d) Aged-SMAT; the scar sizes reduce after SMAT processing in both microstructures.

In Aged samples, there is significant decrease in wear rate after SMAT, which causes the Aged-SMAT to have COF value similar to that of STQ and STQ-SMAT. This corroborates the wear track images in Fig. 13c and d. The extent of SMAT-induced hardening in Aged material is much lower due to precipitate coarsening and possible dissolution of precipitates partially counteracting the hardening of the β-matrix, as discussed above. This likely reduces the material removal rate in the Aged-SMAT thereby generating less debris at the interface. Hardening of the β-matrix can also make it prone to fracture. However, the failure in this case is dominated by the α-β interfaces, which have large strain incompatibility associated with them due to higher strength of α-phase. These interfaces are preferred regions for crack initiation in Aged. After SMAT, the α-β interfaces are fewer due to the coarser microstructure and possible dissolution of α-precipitates. This augments surface ductility resulting in less material removal and lower coefficient of friction. Harder surfaces typically exhibit better wear resistance, although deviation from this can also be found in literature, which includes reports on titanium alloys as well [24,47,48]. The wear rate observed here were in sharp contrast to the trend in hardness, which was as follows: Hv(STQ_SMAT) ≈ Hv(Aged-SMAT) < Hv(Aged) < < Hv(STQ). Although the effect of SMAT on hardness is more profound in STQ than in Aged, the reverse trend is observed for the fretting wear response. Moreover, the wear rate of both Aged as well as Aged-SMAT remains higher than STQ sample, which has the least hardness. This underscores the importance of ductility on the wear behavior. Aside from the mechanical factors, corrosion response of a material also plays a vital role in the fretting wear process, as it involves tribocorrosive environment.

wear scar. The average wear volume trends as follows: VSTQ-SMAT < VSTQ < VAged-SMAT < VAged (Table 3). This is corroborated by the wear scar profiles seen in Fig. 13. Recently, we reported on the wear characteristics of this alloy in STQ and Aged conditions [21]. The higher wear rate for Aged is ascribed to the increased corrosion rate that facilitates formation of wear debris compared to STQ. The presence of wear debris at the fretting interface between alloy and the counterpart eventually results in a three-body wear. Since these are more brittle than the STQ sample, wear becomes more severe in Aged. Herein, we focus on the effect of SMAT. The hardened surface layer in STQ-SMAT sample is associated with more deformation resistance during sliding than that in STQ. Therefore, the wear volume was reduced after SMAT. This is also evident from smaller size of the wear scar of STQ-SMAT in Fig. 13b than that of STQ (Fig. 13a). However, the higher hardness of the STQ-SMAT also indicates a lesser ability to strain harden, that is, less capacity for plastic deformation and lower ductility. Although overall deformation of STQ-SMAT appears less than the STQ, the lower ductility of the former also can enhance the possibility of higher damage of materials [45,46]. As a result, the formation of the wear debris leads to three-body wear in STQ-SMAT, which involves abrasion between the sample and debris. In STQ, softer debris at the interface can reduce the extent of abrasion compared to STQ-SMAT. These factors, therefore, diminish the efficiency of SMAT-induced hardening in reducing the wear rate to some extent. This can also explain the similar friction coefficient values between STQ and STQSMAT samples. 9

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attachment across all the samples at day-1 and appear to have increased by days 4 and 7. The cells were fully confluent at day-7. These results indicate that the surface treatment did not affect the osteoblast attachment and proliferation in this study. Some studies have reported that SMAT leads to an improvement in cell growth and/or differentiation on different materials [12,23,41] and the results have ascribed to changes in surface properties. The results here suggest that measured cell response is unaffected and the alloys continue to exhibit good cytocompatibility after SMAT that is similar to that of cp-Ti. Surface energy is believed to play a key role in driving the cellular response to biomaterials. Processing can alter the microstructure and thereby the surface characteristics to influence the cell response [49]. The surface energy was found to be similar for all the conditions (Table 1), it can be inferred from the present study, that the SMAT did not affect the surface energy apart from a minor change in the roughness of the surface. Taken together, the results of this work show that SMAT does not compromise the cellular response to the alloy, it can be a leveraged for improving the mechanical and other functional properties for biomedical applications. The importance of initial microstructural prior to SMAT processing also was underscored here as different responses were found in single β- and aged β + α microstructure.

Fig. 14. Absorbance values (as a measure of cell viability) for the Ti-Nb-Ta-O alloy in four different condition shown along with cp-Ti. There are no statistically significant differences (p > 0.05) at any given time point between the different samples.

4. Conclusion The conclusions from the present study are as follows:

Interestingly, the wear rate in the present case is in good agreement with the corrosion rate for the different conditions. Reduction in corrosion rates after SMAT-processed samples thus improves the wear resistance for both STQ and Aged.

(1) SMAT of the fully β structured alloy resulted in the formation of deformation substructure on the surface leading to grain fragmentation by generating numerous micro-bands. This resulted in substantial increase of the hardness of the surface. (2) Aged alloy with fine β + α microstructure showed negligible hardening after SMAT as the hardening of β-matrix was counteracted by the softening due to deformation-induced coarsening of the precipitates. (3) The corrosion resistance of both fully β as well as β + α alloys improved after SMAT, which has been attributed to the formation of a denser and stronger passive oxide layer aided by high defect density on the surface. Additionally, in the aged alloy, coarsening of precipitates resulted in reduced number of α-β interface boundaries, thus lowering the propensity of galvanic corrosion between the two phases.

3.7. Biological response Fig. 14 compiles absorbance values from WST-1 assay of osteoblasts cultured on the different samples and is taken as measure of viable cells. At day-1 post-seeding, the absorbance levels on all the samples are similar, which indicates that the cell attachment is similar across all samples. After day-4, the absorbance increases from day-1 due to proliferation of the cells and the number increases further till day-7. Representative fluorescence micrographs for the different processing history and different days are presented in Fig. 15. Actin filaments are stained in green whereas the nuclei are blue. Cells are uniformly

Fig. 15. Representative fluorescence micrographs showing increase in cell numbers with increase in time points for all samples (scale bar = 100 μm). 10

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(4) Fretting wear rate in corrosive environment was found to improve after SMAT in both β and α + β conditions, primarily due to hardened surface layers. Moreover, in the aged alloy, reduction in α + β interfaces possibly improved the ductility at the surface and aided by lower corrosion resistance reduced both wear volume as well as COF. (5) Osteoblast attachment and proliferation on the Ti-Nb-Ta-O alloy did not change after SMAT, possibly because of overall surface wettability remaining unaffected by this process. Taken together, this work shows that SMAT can make the β Ti-Nb-Ta-O alloy more versatile in terms of application in biomedical implant parts.

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CRediT authorship contribution statement Srijan Acharya: Conceptualization, Investigation, Writing - original draft, Writing - review & editing. Arpana Gopi Panicker: Investigation, Writing - review & editing. Vasanth Gopal: Investigation, Writing review & editing. Shaurya Singh Dabas: Investigation, Writing - review & editing. Geetha Manivasagam: Investigation, Writing - review & editing. Satyam Suwas: Conceptualization, Writing - review & editing, Supervision. Kaushik Chatterjee: Conceptualization, Writing review & editing, Supervision. Declaration of competing interest The authors have no conflict of interest to declare. A statement has been added at the end of the manuscript. Acknowledgement The authors acknowledge funding from Uchhatar Avishkar Yojana (Project number IISc_004) supported by the Ministry of Human Resource Development, Govt. of India and the Indian Council of Medical Research. We are grateful to Dr. Amit Bhattacharjee, Scientist and Head of Titanium Alloy Group, Defence Metallurgical Research Laboratory, Hyderabad India for many useful discussions. References [1] Q. Chen, G.A. Thouas, Metallic implant biomaterials, Materials Science and Engineering: R: Reports 87 (2015) 1–57. [2] H. Ikehata, N. Nagasako, T. Furuta, A. Fukumoto, K. Miwa, T. Saito, First-principles calculations for development of low elastic modulus Ti alloys, Phys. Rev. B 70 (17) (2004) 174113. [3] D. Raabe, B. Sander, M. Friák, D. Ma, J. Neugebauer, Theory-guided bottom-up design of β-titanium alloys as biomaterials based on first principles calculations: theory and experiments, Acta Mater. 55 (13) (2007) 4475–4487. [4] E. Eisenbarth, D. Velten, M. Müller, R. Thull, J. Breme, Biocompatibility of β-stabilizing elements of titanium alloys, Biomaterials 25 (26) (2004) 5705–5713. [5] S. Acharya, A.G. Panicker, D.V. Laxmi, S. Suwas, K. Chatterjee, Study of the influence of Zr on the mechanical properties and functional response of Ti-Nb-ta-Zr-O alloy for orthopedic applications, Mater. Des. 164 (2019) 107555. [6] S. Acharya, P. Gupta, K. Chatterjee, S. Suwas, Microstructure, texture and mechanical properties after cold working and annealing in a biomedical Ti-Nb-Ta alloy, Materials Science Forum, Trans Tech Publ, 2018, pp. 2465–2470. [7] T. Saito, T. Furuta, J.-H. Hwang, S. Kuramoto, K. Nishino, N. Suzuki, R. Chen, A. Yamada, K. Ito, Y. Seno, Multifunctional alloys obtained via a dislocation-free plastic deformation mechanism, Science 300 (5618) (2003) 464–467. [8] D. Gordin, R. Ion, C. Vasilescu, S. Drob, A. Cimpean, T. Gloriant, Potentiality of the “gum metal” titanium-based alloy for biomedical applications, Mater. Sci. Eng. C 44 (2014) 362–370. [9] C.A. St Pierre, M. Chan, Y. Iwakura, D.C. Ayers, E.A. Kurt-Jones, R.W. Finberg, Periprosthetic osteolysis: characterizing the innate immune response to titanium wear-particles, J. Orthop. Res. 28 (11) (2010) 1418–1424. [10] M. Lai, K. Cai, Y. Hu, X. Yang, Q. Liu, Regulation of the behaviors of mesenchymal stem cells by surface nanostructured titanium, Colloids Surf. B: Biointerfaces 97 (2012) 211–220. [11] S. Bahl, P. Shreyas, M. Trishul, S. Suwas, K. Chatterjee, Enhancing the mechanical and biological performance of a metallic biomaterial for orthopedic applications through changes in the surface oxide layer by nanocrystalline surface modification, Nanoscale 7 (17) (2015) 7704–7716. [12] S. Bahl, B.T. Aleti, S. Suwas, K. Chatterjee, Surface nanostructuring of titanium imparts multifunctional properties for orthopedic and cardiovascular applications, Mater. Des. 144 (2018) 169–181.

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