Enhancing Passivation Behaviors and Wear Resistance of Biomedical Ti-15Mo Alloy via {332} Twinning Pre-Tension and Aging

Abstract

The purpose of this investigation was to determine the extent to which a combination of pre-deformation and aging procedures could enhance the corrosion and wear resistance of the Ti-15Mo alloy for biomedical materials. The initial grains were refined with {332} twins and other defects after the pre-deformation, increasing the amount of precipitation nucleation and forming finer α phases during aging. The finer precipitates with numerous boundaries are beneficial in forming a thicker passivation film in phosphate-buffered saline (PBS) solution, and the corrosion resistance of the 20% pre-deform plus aged alloy is improved by 4.23 times. The result indicates that the corrosion passivation behavior and film structure of the biomedical Ti-15Mo alloy are significantly influenced by its microstructure. The worn track and debris of the alloy after pre-tension and aging gradually decreases with the increase in pre-deformation rates, caused by the increase in the hardness and wear resistance. The mechanism of corrosion and wear resistance in PBS solution were revealed, which showed the pre-tension and aging processes that were performed on the Ti-15Mo alloy. This study proposes that pre-tension with {332} twins contributed to precipitation refinement, which would enhance the passivation behaviors and wear resistance of the biomedical β titanium alloy.

1. Introduction

Titanium (Ti) alloy with excellent mechanical properties and biocompatibility has been widely used as a biomaterial in the past few decades [1,2,3]. The increase in human lifespan and a fast lifestyle has caused a higher requirement for biomedical implants that are more durable and reliable. However, traditional Ti alloys have a nonnegligible risk of toxicant elements due to corrosion behaviors, frictional wear and ion dissolution during long-term service. Thus, biomedical β-Ti alloys are designed to meet the non-toxicity requirements [4] and possess superior biological performance and compatibilities in the human body environment. On the other hand, the loss and failure of biomaterials are noteworthy issues due to long-term corrosion and friction during the service process. Excellent corrosion and wear resistance is also desired and favorable for extending the service life of biomedical implants, so that Ti implants should endure complex corrosion and wear in vivo [5].

As is well known, the crucial strategy for performance improvement in metals is microstructural refinement. In a previous study by the authors, the improved microstructure and strength were found in the Ti-15Mo-3Al-2.7Nb-0.2Si alloy after cold pre-rolling and aging [6]. The Ti-Nb alloy after pre-deformation and subsequent aging also presented the optimized β matrix texture and morphology of the α precipitates [7]. Similarly, the Ti29Nb13Ta4.6Zr alloy with nanoscale α precipitates after cold rolling and aging demonstrated improved mechanical properties and decreased the passivation current density [8]. Ji et al. reported that the microstructure of the Ti alloy significantly influenced the corrosion resistance via improving the nucleation and growth of the passive film [9]. This passive film acts as a barrier, preventing biological and chemical reactions between biomedical implants and the corrosion environment. The microstructural refinement facilitates the formation of a dense passive film and improves the corrosion resistance [10]. Enhanced passive films in PBS solution can be obtained by refining α precipitates in β-Ti alloys. The content of hypervalent oxides in the passive film also increases with the refinement of the microstructure [11]. A biomedical Ti-15Mo alloy does not have toxicant elements compared with other Ti alloys, and is easy to deform with a {332} <113> twinning mechanism [12,13,14]. The twins are different from slipping dislocations and have special effects on the precipitation behavior during the aging process after pre-deformation. However, pre-deformation with the twinning mechanism and the precipitation behaviors of β-Ti alloys are scarcely reported. Meanwhile, the passivation behaviors and corrosion resistance of the Ti-15Mo alloy after pre-deformation and aging should be investigated systematically.

In order to obtain excellent corrosion and wear resistance, this study introduced {332} twins and other defects via pre-tension, refining the coarse grains and providing assistant nucleation for the aging precipitation of the Ti-15Mo alloy. In contrast to conventional solution and aging treatments [15,16], the alloy after pre-deformation and aging treatments is desired for extending the service life of biomedical applications. Microstructural evolution and its effects on the corrosion resistance were also investigated using electron backscatter diffraction (EBSD), X-ray diffraction (XRD), transmission electron microscopy (TEM), electrochemical impedance spectroscopy (EIS) and X-ray photoelectron spectroscopy (XPS). This study would more deeply reveal the mechanism of corrosion and wear resistance of the alloy after pre-tension and aging.

2. Experimental Materials and Methods

2.1. Specimen Preparation

The Ti-15Mo alloy used in this study was melted as an industrial ingot by vacuum arc melting three times. The chemical composition is listed as follows: 14.7 mass% Mo, 0.08 mass% O and Ti balanced. The transformation temperature of the α/β phase is approximately 755 °C via the metallography method. The billet of the alloy was hot-rolled to a Φ24 mm bar followed by homogenization and forging. After the solution treatment (ST) at 790 °C for 1 h and quenching, the pre-tension of the alloy was conducted on an MTS 370.02 universal testing machine (Eden City, MN, USA) at room temperature under 0.001 s⁻¹. The samples after the 3%, 9%, 15% and 20% deformation rate were marked as Tens-3, Tens-9, Tens-15 and Tens-20, respectively. According to our previous studies [17,18], the aging treatment following ST/pre-deformation was conducted at 550 °C for 8 h. The Tens-3, Tens-9, Tens-15 and Tens-20 specimens after aging treatment were marked as Tens-3-A, Tens-9-A, Tens-15-A and Tens-20-A, respectively. After pre-tension and aging, the alloys were, respectively, subjected to various characteristic tests, including microstructure, hardness, corrosion and wear resistance, as shown in Figure 1. The specimens for the tests were cut by DF7745 electrical discharge machining (Taizhou, China), and paralleled to the tensile axis.

2.2. Electrochemical Measurements

The corrosion test was conducted in electrolyte using the ivium electrochemical workstation (Eindhoven, Holland). The commercial PBS solution with 7.4 PH was chosen in this study with the following composition (unit: mM): 139.7 Cl⁻, 153 Na⁺, 8 HPO₄²⁻, 4.7 K⁺ and 2 H₂PO₄⁻, which was manufactured by Labcoms^TM LIFE SCIENCES (Shanghai, China). The electrochemical measurement was performed with the conventional three-electrode cell system, including a working electrode, a platinum sheet as the counter electrode and a reference electrode of saturated AgCl. The EIS, potential dynamic (PD) polarization and Mott–Schottky (M-S) method were employed in this study. The open-circuit potential (OCP) was recorded for 3600 s to ensure stability before other electrochemical tests. The negative voltage of −1.5 V was loaded for 5 min before the OCP tests to dissolve the oxide layer formed in the atmospheric environment during the sample preparation. The EIS experiment was performed with a scan frequency range from 10⁻² Hz to 10⁵ Hz. And the EIS data were fitted to an equivalent electric circuit (EEC) model by the ZSimpWin software, which is integrated in the electrochemical workstation. The M-S test was conducted in a potential range from −1000 mV to 500 mV. The PD polarization curve was recorded from −1000 mV to 2000 mV at a scanning rate of 1 mV/s. For each condition, at least two specimens were tested to ensure reproducibility.

2.3. Friction Tests and Characterization Methods

The friction test was conducted at an ambient temperature using Bruker’s universal mechanical tester (UMT, Billerica, MA, USA). Prior to testing, the specimen for the friction tests underwent mechanical grinding and polishing to achieve a mirror-like finish surface. The specimens were immersed in PBS solution as the friction medium during testing. The applied load was 1 N, the frequency was set at 1 Hz and a stroke length of 2 mm was used. The duration of the test was 30 min. An aluminum oxide ball with a diameter of 6 mm was employed as the friction pair. The worn tracks of tribological experiments were observed using a DCM8 confocal microscope (Wetzlar, Germany). To ensure reliability of results, three replicate wear tests were conducted for each specimen under identical conditions.

The hardness of the alloy was measured by a micro-Vickers hardness tester (HV-1000, Xi’an, China), with 500 g for 5 s and measured 10 times to obtain an average value. The JSM6460 scanning electron microscope (SEM, Tokyo, Japan) was utilized to observe the morphology of the microstructure, worn tracks and debris. Electron backscatter diffraction (EBSD) analysis was conducted using an Oxford Instruments AZtecSynergy operated at 20 kV and a step of 0.3 μm. The EBSD data were post-processed by the channel 5 software. The morphology of deformed structures and α precipitates were observed using a JEM-3010 transmission electron microscope (TEM, Tokyo, Japan) operated at 100–300 kV.

3. Results

3.1. Microstructure of ST and Its Evolution during Pre-Tension

Figure 2 illustrates the orientation imaging map (OIM) and pole figures (PFs) of the Ti-15Mo alloy before pre-tension. The size of the grains in the ST alloy is about 50 μm and presents a regular shape (Figure 2a). It suggests that recrystallization had completed after the ST at 790 °C for 1 h. The ST specimen consists of a single β phase, which has a body-centered cubic (BCC) crystal structure. The initial texture of the alloy investigated by EBSD (Figure 2) and the OIM of the rolling direction (RD) illustrated that no dominant texture had been formed after the rolling and ST processes. The distribution of the main orientations in the Ti-15Mo alloy after the ST is relatively random. However, a slight strong {100} orientation could be observed from the PF analysis in Figure 2b. This recrystallized texture was liable to form in bcc metals [19,20,21], and also indicated that the recrystallization of the ST alloy had completed. No twin was observed in Figure 2a and it indicated that the stacking fault energy (SFE) of the Ti-15Mo alloy was insufficiently low to form an annealing twin during the ST and water quenching processes. Differences in the ω and α″ phase were also not observed in the alloy due to the low resolution of EBSD. However, previous works of the authors had investigated the ST Ti-15Mo alloy with TEM and revealed that these metastable phase transformations had taken place slightly during quenching [17,18].

{332} <113> twins of β-Ti alloys could be formed during cold deformation and had 50.57° misorientation [22,23,24,25]. The pre-tension in this work was performed in the Ti-15Mo alloy at the 3%, 9%, 15% and 20% rate at room temperature, respectively. The OIMs and boundary character distribution (BCD) of the EBSD after pre-tension are shown in Figure 3. The boundaries with 10~49° misorientation are marked with green color. Black color marks the boundaries where the misorientation was over 53°. The 49~53° misorientation boundaries are marked with red color, and most of them should be {332} <113> twins formed during the pre-tension deformation of the Ti-15Mo alloy. Figure 3a shows the OIM and BCD of the Ti-15Mo alloy, which was performed by 3% pre-tension deformation. The density difference in the {332} <113> twins is significantly observed in different recrystallized grains due to the original texture. {332} <113> twins are apt to forming in some certain texture matrices, which was decided with the difference in the critical resolved shear stress (CRSS) during deformation [13]. The grain with a higher Schmidt factor (SF) would be apt to form {332} <113> twins and deform with the twinning mechanism [26]. It could be observed that a few {332} <113> twins appeared in the interior of recrystallization grains. This phenomenon was also reported by Wu et al. while {112} twins were investigated [27]. The Ti-15Mo alloy with 3% pre-tension is considered in an earlier stage for tension deformation, so some {332} <113> twins terminated in the interior of grains. With an increase in pre-tension deformation, some twins propagated to the neighboring grains across the low-angle boundary (LAB), but some ended at the high-angle boundary (HAB) and induced a new twinning formation in the neighboring grain due to the large stress concentration. Figure 3b shows the OIM and BCD of the alloy with 9% pre-tension and presents a significant difference from Figure 3a. Firstly, the number of twins increased obviously with the increase in the pre-tension deformation. Boundaries crossing the black line in Figure 3b were still about 50.57° misorientation and suggested that {332} <113> twins were the main deformation mechanism while the alloy was conducted by 9% pre-tension. Secondly, most of the twins have propagated to the whole grain, indicating that the length of twins increased with the development of pre-tension deformation. However, the width of the twins in the Ti-15Mo alloy with 9% pre-tension never has an obvious change and indicates that the amount and length of twins increased with the increase in pre-tension deformation but the width did not.

The OIM and misorientation of boundaries in the alloy with 15% pre-tension are shown in Figure 3c. In contrast to the 9% pre-tension, Figure 3c presents a higher density of twins in the Ti-15Mo alloy after 15% pre-tension deformation. Furthermore, lots of {332} <113> twins had different orientations and were distributed crosswise in the initial grain. The sufficient stress caused by the increased pre-tension deformation is considered to induce these crisscrossed twins due to the requirement for a higher CRSS. The crisscrossed twins also refined initial grains and formed lots of sub-structures with different orientations. Some misorientation of twin boundaries at the black line in Figure 3c is lower or higher than 50.57° obviously and indicates that the sub-structures might be rotated and form deviation from the {332} <113> twinning misorientation during 15% pre-tension deformation. Even so, the interface of the twin and matrix still presents a rough straight shape due to the sub-structures maintaining the morphology of twinning boundaries. The alloy performed by 20% pre-tension presents a different microstructure in contrast to the 3%~15% deformation. Twinning boundaries formed in high SF grains conducted sufficient deformation and rotation, which induced some grains to be refined due to there being lots of irregular sub-structures (in Figure 3d). Although boundaries of these sub-structures at the black line in Figure 3d are no longer straight, most of them are around 50.57° misorientation and show as the HABs. It is different from the Ti-15Mo-3Al-2.7Nb-0.2Si alloy, which is dominated by dislocation motion during cold-rolled deformation and showed low-angle grain boundaries [6]. Another interesting result could be obtained from the fact that twins are rarely found in the <001> orientation grains (red color in Figure 3) compared with other orientation grains. According to the previous literature [28,29], these <001> orientation grains are not easy to form {332} <113> twins during pre-tension deformation due to their low SF and requirement for a higher CRSS [28,30].

Figure 4 shows the boundary size distribution (BSD) of the alloy after different pre-tensions. As the deformation increased, the initial grains were continuously refined by various twin variants. Compared with the ST specimen, the alloy after pre-deformation has more refined sub-structures. The size of sub-grains in the Tens-3 specimen is 0.5~4 μm, even over a 46.3% proportion of them are approximately 1.0 μm, which is caused by twinning formation during the pre-tension. Over a 53.1% proportion of the sub-grains have a size of approximately 1.0 μm, exhibiting more refined sub-structures in the alloy after 9% pre-tension. The number of sub-grains in the Tens-15 specimen is slightly higher than that of Tens-9. The number of grain sizes of about 0.5 μm became maximum in the Tens-20 specimen, due to the formation of nanoscale sub-structures at the twin and twin-crossing positions. These deformed sub-grains in the alloy after pre-tension would have an influence on precipitation behaviors during the subsequent aging.

As seen in Figure 5 in the TEM images, the microstructure of the alloys subjected to 9% and 20% pre-tension is shown. The bright-field (BF) image in Figure 5a was used to analyze the morphology of the twins and boundaries of the Ti-15Mo alloy deformed with 9% pre-tension. The width of the twins is less than 1 μm and the boundaries present as a straight shape. The TEM of the 9% pre-tension alloy also shows that the orientations of the twins and the matrix have a significant difference with the white and black contrast. Figure 5b presents the selected area electron diffraction (SAED) pattern of the twins (a red circle in Figure 5a) and indicates their characteristic of {332} <113> twins. Furthermore, in the alloy after the Tens-9 process, the athermal ω phase has been identified by the diffraction spots at 1/3 and 2/3 [211] M (in Figure 5b). A different morphology is shown in Figure 5c and indicates that evolution of the twinning should take place during pre-tension from 9% to 20%. The BF image of TEM (in Figure 5c) illustrates that dense and crossed twins have cut the original grain into many small sub-structures. Furthermore, the density of dislocations in these twins and sub-structures is significantly higher in contrast to that of the 9% pre-tension alloy (in Figure 5a). The SEAD pattern in the bottom-left corner in Figure 5c also indicates that many fine ω phases formed in the alloy after 20% pre-tension. These ω phases would serve to reduce the energy barrier and play a crucial role in assisting the nucleation of α precipitates during aging.

Figure 6 shows the synthetic images of the {332} twin boundary and Kernel average misorientation (KAM) of the alloy after 3%, 9%, 15% and 20% pre-tension, respectively. The boundaries of the twins are exhibited as brown lines, which are automatically generated by channel 5 software with the corresponding misorientation between the adjacent regions. In this study, the 49–53° misorientation is identified as {332} twin boundaries, because the misorientation of {332} twins induced by pre-tension could have a small deviation from the standard of 50.57°. Figure 6a shows that a few twin boundaries appear in some initial grains of the alloy after 3% pre-tension deformation. The alloy after 9% pre-tension formed more twin boundaries, which could be observed in all grains (in Figure 6b). With the increase in pre-tensive deformation, these twinning variants with different orientations exhibited a cross-division observation, resulting in the refinement of the initial grains. While the pre-tensive deformation was over 9%, misorientation of some boundaries was beyond the range of 49–53°. This deviation from the twinning characteristic might be caused by the torsion of {332} twins. The further microstructural refinement of the Ti-15Mo alloy would occur with the increase in the cross-division and torsion of the twins. These sub-structures became smaller and more boundaries deviated from the standard misorientation of {332} twins (shown in Figure 6c,d). The more deviated boundaries were observed with the increase in the pre-tension (in Figure 6a–d), indicating that more twin structures occur due to torsion with the further deformation. These boundaries still play an important role in the microstructural refinement, even if they no longer have a perfect {332} twin boundary.

The KAM maps were used to analyze the corresponding lattice rotation after pre-tension. In this study, the red color in the KAM map denotes the region with a 5° rotation, signifying the highest lattice rotation, while the blue color indicates a negligible rotation with 0° value. The KAM maps of the Tens-3, Tens-9, Tens-15 and Tens-20 specimens are presented in Figure 6. The KAM of the twins displays a higher value compared to the surrounding matrix, indicating a greater rotation and more defects due to deformation. Notably, the KAM value in the joint region of the matrix and twins is higher than that of the surrounding regions, indicating that the twin boundaries are prone to acting as dislocation sources during pre-tensive deformation. The difference in the KAM from Figure 6 also illustrated that more dislocations and residual stresses were formed in the alloy with the increase in the pre-tensive deformation.

3.2. Microstructure and Precipitation Behaviors during Aging

Generally, the α phase with a close-packed hexagonal (CPH) crystal structure will precipitate from the β phase during aging in the STed Ti-15Mo alloy. In this study, the alloy after the solution and aging treatments is marked as ST-A for short. Similarly, the 3%, 9%, 15% and 20% pre-deformed alloy after aging is labeled as Tens-3-A, Tens-9-A, Tens-15-A and Tens-20-A, respectively. Figure 7 presents the SEM and TEM morphology of α precipitates in the ST-A sample. The ST over T_β causes the β matrix phase to be metastable, and more α precipitates would form during aging to change the stability of the alloy. However, the coarse grains in the ST alloy provide sufficient spaces for the growth of the α precipitates, resulting in a relatively large size of the α phase (in Figure 7). The TEM image of the ST-A sample is shown in Figure 7b and illustrates that the size of α precipitates in the ST-A sample is approximately 50 nm in width and 5 μm in length. The result of SAED, shown in the top-right corner of Figure 7b, exhibits that there are only α precipitates after the 550 °C aging treatment.

Figure 8a–d show the SEM microstructure of the Tens-3-A, Tens-9-A, Tens-15-A and Tens-20-A specimens, respectively. In contrast to the morphology of the α precipitates in the ST-A specimen (in Figure 7), the α precipitates formed during aging in the pre-deformed alloy are significantly finer. The pre-tension influenced the precipitation behaviors due to two main factors. Firstly, the dislocations and ω metastable phases formed during pre-tension facilitate the nucleation of the β→α phase transformation in a non-equilibrium condition. Secondly, the {332} twins induced by pre-tension refined the initial grains, restricting the growth of α precipitates and preventing the formation of coarse precipitation. It can be observed from Figure 8 that the size of the α precipitates formed during aging becomes finer and distributes more randomly with the increase in the pre-deformation. In the pre-tensive deformation over 15%, the length of precipitation is less than 1 μm (in Figure 8c,d), which is significantly refined compared with the size of the α precipitates in the ST-A specimen. The nanoscale α precipitates formed during aging in the pre-tensive alloy illustrate short-needle morphology, and distribute in the β matrix grains homogeneously. Figure 8e–h show the TEM microstructure of the alloy after 3%, 9%, 15% and 20% pre-deformation and aging. Different from the ST-A specimen, two changes could be observed in the alloy after pre-deformation and aging. Firstly, numerous dislocations and ω metastable phases formed during pre-tension provide sufficient non-equilibrium nucleation sites for the precipitation phase transformation. Secondly, the alloy after pre-tensive deformation formed numerous twins and sub-structures, which restricted the growth of α precipitates due to insufficient space, resulting in microstructural refinement. The amount of α precipitates increased, the size of them decreased and the distribution became more random with the increase in the pre-deformation rate (as seen in Figure 8).

3.3. Passivation Behaviors and Corrosion Resistance of Biomedical Ti-15Mo Alloy

Figure 9a shows the open-circuit potential (OCP) curves of the ST-A, Tens-3-A, Tens-9-A, Tens-15-A and Tens-20-A specimens in PBS solution. It was observed that as corrosion proceeded, the corrosion potential in all the specimens increased, indicating that the surface of the Ti-15Mo alloy spontaneously formed an oxide film, which played a role in inhibiting corrosion. The formation of the oxide film on the surface of titanium metal affects the electrochemical reaction at the corrosion electrode, causing the open-circuit potential to shift positively, which is a gradual passivation process. As the metal is immersed in PBS solution for a longer period, the thickness of the passivation film increases, enabling protection against the continuous dissolution of the metal. The initial stage of the OCP curve shows a rapid increase in the positive potential, and as corrosion proceeded, the change in the OCP becomes smaller, eventually reaching a dynamic equilibrium. The result indicates that the stabilized structure of the passivation film (chemical composition, thickness, defect concentration, etc.) has been formed completely. The required time for reaching a stable OCP reflects the difference in the passivation speed of the Ti-15Mo alloy after different pre-tension and aging. Compared with the ST-A sample, in general, the alloy after pre-tension and aging needs a shorter time to be a stable OCP. Furthermore, this trend is more obvious with the increase in the pre-deformation rates of the Ti-15Mo alloy. Furthermore, the open-circuit potential after 3600 s became stable in all specimens, and the value increased with the increase in the pre-tensive deformation (in Figure 9b), indicating an enhanced corrosion resistance of the alloy.

Figure 10a,b present the Nyquist plots and the Bode plots of the alloy after pre-tension and aging. The results indicate the noticeable difference in the semi-circle radius of the Nyquist plots due to different pre-deformations conducted in the ST-A, Tens-3-A, Tens-9-A, Tens-15-A and Tens-20-A specimens. In principle, a larger circle radius in the Nyquist plots suggests a higher impedance of the passivation film. Notably, the passivation film on the ST-A specimen with a smaller radius was attributed to its initial coarse grains and rod-like precipitates. The Nyquist plots of the passivation films on all the specimens shown in Figure 10a illustrate an incomplete single capacitor arc, indicating their typical capacitance characteristics. The size of the impedance arc reflects the electrochemical reaction rate of the passivation film of the alloy. Generally, a larger diameter of the impedance arc suggests a slower charge transfer process and a lower corrosive rate. Compared to the ST-A alloy, the passivation film on the pre-deformed and aged alloy exhibits a significantly larger impedance arc diameter, slowing down the corrosion reaction effectively. This observation is caused by pre-tension, which significantly refined the initial grains and facilitated the formation of finer precipitates. The more ultra-fine microstructures of the pre-deformed and aged alloys possess massive boundaries and promote the formation of more stable passivation films. The impedance arc of the Tens-3-A specimen was slightly larger compared with other pre-tension and aged specimens, due to the difference in the vacancy concentration in the passivation film.

The Bode plots of the ST-A, Tens-3-A, Tens-9-A, Tens-15-A and Tens-20-A specimens are shown in Figure 10b. The results reveal a single peak in the Bode-phase plots for each specimen with a range of 50~75°. The observation indicates that the passivation film formed during OCP has a single time constant, resulting in relatively stable passivation behaviors in PBS solution. In both the Bode and impedance plots, two frequency bands can be observed. In a high-frequency region (10³~10⁵ Hz), the phase angle values dropped to 0°, and the impedance almost remains constant, which is a reflection of the resistance of the electrolyte between the reference and working electrodes. However, in the low-frequency region (10⁻²~10² Hz), the slope of impedance/frequency is close to −1, suggesting the capacitive behavior of passivation films on all specimens. This observation also suggests that a stable oxide film has formed on the alloy’s surface, demonstrating the stability of the surface–electrolyte interaction. Additionally, the phase angle decreases with the increase in frequency within the range of 10 to 10³ Hz. The results of different frequency regions suggest that there are different electrochemical mechanisms due to the characteristics of the passivation film. In the high-frequency region, the phase angle was close to 0, reflecting that the passivation film effectively inhibited the impedance change.

The stable passivation film improves the corrosion resistance of the Ti-15Mo alloy in PBS solution and improves its biomedical performance. The result of Figure 10b reveals that all specimens exhibit a single time constant. To evaluate the resistance of the passivation film on the ST-A, Tens-3-A, Tens-9-A, Tens-15-A and Tens-20-A specimens in PBS solution, an electrical equivalent circuit (EEC) was used to analyze the EIS data. The values of the EEC elements are calculated via the experimental EIS fitting results using ZSimpWin software and are shown in the top-right corner of Figure 10a and Table 1. The elements of the EEC have specific meanings: R_s signifies the solution resistance, R_f is the passivation film resistance and CPE stands for the constant phase element. Noteworthy, the CPE is used instead of a traditional capacitor due to the phase angle values being less than 90° in the low-frequency range (10⁻²~10² Hz), suggesting a nonideal capacitor at the electrode/electrolyte interface [31]. This deviation from ideal capacitive behavior may be caused by varying relaxation times due to microscopic inconsistencies in the passivation film. The impedance of a CPE is expressed as follows: [32]

(1)$z_{CPE}={\displaystyle \frac{1}{{\left(j\omega \right)}^{n}CP{E}_f}}$

where Z_CPE signifies the constant of CPE, j as the imaginary number and ω as the angular frequency. The value of n determines whether the constant phase element behaves like a capacitor, resistor or inductor. For n = 1, it functions like an ideal capacitor, for n = 0, as a pure resistor and for n = −1, as an inductor. The n value of all specimens in this study is above 0.82, according to Table 1, indicating that the passivation film on the Ti-15Mo alloy’s surface behaves more like a capacitor. The R_f and CPE_f, as the elements of the EEC, are listed in Table 1, representing the resistive and capacitive responses of the passivation film. The effective capacitance (C_eff) of the passivation film can be calculated using the following equation [33,34,35]:

(2)$C_{eff}={\left({R_f}^{1-n}{CPE}_f\right)}^{{\displaystyle \frac{1}{n}}}$

Since the passivation film shows similar capacitor characteristics, its thickness can be likened to the space between parallel plates in a capacitor. Consequently, the passivation film’s thickness is calculated as the following equation [33]:

(3)$L_{SS}=\left({\epsilon }_r{\epsilon }_0\right)/C$

where ε₀ is the vacuum permittivity and 8.854 × 10⁻¹⁴ F/cm, as is commonly known [34]. ε_r represents the permittivity of the Ti-based oxide film, which has been identified in the previous literature as 65 [33].

Table 1 presents the EEC fitted parameters including the calculated C_eff, L_ss values for the specimens. It can be observed that the R_s values of all specimens in PBS solution are quite similar. R_f shifts from 123 to 843 kΩ·cm² due to variations in the passivation films. In PBS solution, the CPE values were 6.53 μΩ⁻¹·cm⁻²·sⁿ for the ST-A specimen and 2.43 μΩ⁻¹·cm⁻²·sⁿ for the Tens-20-A specimen, suggesting a significant difference in capacitor characteristics due to the microstructural refinement. The thickness (L_ss) of the passivation film in the alloy after pre-tension and aging is approximately 1.5~4.4 times more than that of the ST-A alloy. The L_ss values calculated using Equation (3) show that the ST-A specimen has a film thickness of 6.32 nm, Tens-3-A has 9.43 nm, Tens-9-A has 15.0 nm, Tens-15-A has 24.12 nm and Tens-20-A has 27.71 nm. These results suggest that the thickness of the passivation film increases with the increase in the pre-deformation. Specifically, the L_ss value of the Tens-20-A specimen is 4.4 times that of the ST-A specimen. The variation in the EEC parameters suggests that the microstructure is significantly influenced by the formation of passivation films on the specimens. The twin boundaries, ω phase and dislocations formed during the pre-deformation refined the initial grain and deduced finer precipitates during aging, which also increases the nucleation sites for forming a thicker passivation film, which would decrease the electrode reaction rate and enhance the corrosion resistance of the metal. The refined grains and finer precipitates are considered as dominant factors, which promoted the formation of Ti- and Mo-based oxides and improved the corrosion resistance.

Figure 11a presents the Mott–Schottky (M-S) plots and fitted curves for the passivation film on the ST-A, Tens-3-A, Tens-9-A, Tens-15-A and Tens-20-A specimens in PBS solution. The M-S curves illustrate linearity within the potential range of 0 mV to 500 mV. In contrast, the nonlinearity of the Mott–Schottky curves was caused by the interference of the Faraday reaction on space charge capacitors like TiO₂ and Ti₂O₃, which was reported in the other metal and solution [36]. In this study, the potential interval of 0 mV to 500 mV (vs. AgCl) was chosen for calculating the donor concentration. In general, M-S measurements reveal that the passivation films on all alloys act as n-type semiconductors similar to other Ti alloys, with oxygen vacancies as the dominant charge carriers [36]. These vacancies introduce additional active sites, influencing the chemical reactions and the surface energy of the passivation film. The oxygen vacancies also facilitate the formation of the passivation film, since the oxygen vacancy formation energy (2.7 eV) is lower than that of a titanium interstitial (4.7 eV). Hence, oxygen vacancies are the primary donors in the passivation film, and their densities can be calculated from the slope of the Mott–Schottky fitted lines, as shown in Table 2. According to the M-S relation, the space charge capacitance of N-type semiconductors can be calculated using the following equation [36,37]:

(4)${\displaystyle \frac{1}{C^2}}={\displaystyle \frac{2}{\epsilon {\epsilon }_{0}e{N}_D}}(E-E_{FB}-{\displaystyle \frac{kT}{e}})$

(5)$N_D={\displaystyle \frac{2}{e{\epsilon }_r{\epsilon }_{o}b}}$

where b represents the slope of the fitted lines and N_D denotes the donor concentration (cm⁻³). Table 2 displays the values of the slope of the M-S fitted lines and the corresponding ND values calculated using Equations (4) and (5). The calculated N_D values for all specimens fall within the range of 6.28 × 10²⁰ to 11.4 × 10²⁰ cm⁻³, which is in a narrow range. The concentration of oxygen vacancies in the passivation films remains relatively constant regardless of pre-deformation. However, a slight variation is still observed in the Tens-3-A specimen, indicating a lower N_D value than in the Tens-9-A and Tens-15-A specimens. This result explains a larger impedance arc in the Nyquist plots observed in the Tens-3-A specimen, though there is a thinner passivation film compared to the Tens-9-A specimen.

Potential dynamic (PD) polarization measurements were conducted to analyze the performance of the passivation films in PBS solution. Figure 11b displays the PD polarization curves of the ST-A, Tens-3-A, Tens-9-A, Tens-15-A and Tens-20-A specimens in PBS solution. Despite the surfaces of the Ti alloy being covered with typical barrier-type passivation films, electronic and ionic conductivity can still generate the current. Figure 11b illustrates two distinct passivation platforms for the current density in the polarization ranges of −300 mV~300 mV and 500 mV~1500 mV (vs. Ag/AgCl), indicating the stability of the passivation films of the Ti-15Mo alloy within these voltage ranges. These films are crucial in preventing direct contact between the electrolyte and fresh metal surface. The values of the corrosion voltage (E_corr), corrosion current (I_corr) and passivation current (I_pass) derived from fitting and analyzing the PD polarization curves are summarized in Table 3. The E_corr values of all specimens are similar. However, in contrast to the ST-A alloy, the I_corr values are significantly lower due to the enhancement in the oxide film for the alloy after pre-tension and aging.

Passivation currents in the first platforms (1st I_pass) decrease with the increase in pre-tension, indicating that the corrosion resistance of the alloy after pre-rolling and aging is significantly improved. The values of 1st I_pass at 0 mV in the ST-A, Tens-3-A, Tens-9-A, Tens-15-A and Tens-20-A specimens are 1.38, 0.42, 0.73, 0.80 and 0.33 μA/cm², respectively. However, the values of 2nd I_pass at 1000 mV do not exhibit this trend. As the potential in the human body is generally lower than 500 mV, giving more attention to the 1st I_pass value is valuable. The results from Table 3 indicate that the corrosion resistance of the Tens-20-A specimen is 4.2 times better than that of the ST-A specimen. The alloys deformed by pre-rolling form lots of {332} twins, ω phases and dislocations, causing the refinement of grains. Subsequently, finer α precipitates and their boundaries provide more sites and channels to form a thicker oxide film, which is a critical factor for the corrosion resistance of the alloy.

3.4. Friction Performance of the Alloy after Pre-Tension and Aging

Figure 12a shows the friction coefficient of the Ti-15Mo alloy in PBS solution, suggesting that a significant fluctuation is observed in all specimens with the duration of friction. This phenomenon was caused by the friction characteristics and friction type of the Ti alloy. The exfoliation of abrasive particles in the PBS solution leads to the observed fluctuation in the friction coefficient. This is consistent with the previous literature [38], which had reported instability of the friction coefficient in Ti alloys. Figure 12b presents the microhardness and wear loss of the ST-A, Tens-3-A, Tens-9-A, Tens-15-A and Tens-20-A specimens after friction in PBS solution. By analyzing the mechanical properties of the aged alloys following different pre-tensions, a relationship between the friction coefficient, wear and microstructure is established. The results affirmed that the increased hardness of the pre-deformed and aged alloy provided a positive impact on the wear resistance. It is an effective method for improving the hardness and the wear resistance of the alloy via microstructural refinement, such as the pre-tension and aging processes.

Figure 12b shows that the hardness values of the aged alloy increase with the increase in the pre-tensive deformation. Compared with the ST-A specimen, the hardness value of the Tens-20-A specimen increases to 335.4 from 272.0, indicating a 23.3% increase. The result is caused by two factors. Firstly, the {332} twins contribute to the strengthening effect of refinement by segmentizing the initial grains. Secondly, the finer nanoscale α precipitates formed during pre-tension cause a severe dispersion in the strengthening effect. Figure 13a,c,e,g,i show the three-dimensional morphology of the worn tracks of the ST-A, Tens-3-A, Tens-9-A, Tens-15-A and Tens-20-A specimens, respectively. The wear loss in PBS solution decreases with the increase in the hardness. Although the passivation film on the specimen influences the friction behaviors, the decisive factor is the hardness. A higher hardness would result in a lower wear loss. The result of the wear loss indicates that a significant enhancement in the wear resistance occurred in the alloy after pre-tension and aging. The wear loss of the ST-A specimen is 3.32 × 10⁷ μm³, but that of the Tens-20-A specimen is 2.24 × 10⁷ μm³ in the same friction condition, exhibiting a 33% increase in wear resistance. Additionally, the Tens-20-A specimen had the smallest friction coefficient (in Figure 12a); thus, the best wear resistance could be concluded as the lowest wear loss.

The SEM images in Figure 13b,d,f,h,j present the micro-morphology of the worn tracks on the ST-A, Tens-3-A, Tens-9-A, Tens-15-A and Tens-20-A specimens, respectively. The results illustrate ridges and grooves along the sliding direction with randomly distributed metallic debris, suggesting abrasive wear as the primary wear mechanism. The worn tracks on the specimens mostly have a flat appearance, suggesting surface damage with micro-cutting and flaky stripping. The ST-A specimen displays deep grooves and larger particles adhering to the surface, suggesting adhesive wear resulting from the continuous formation and attachment of particles during friction (in Figure 13b). Numerous grooves can be seen on the surface along the sliding direction, with a layer of worn debris covering the wear grooves. Plastic deformation of the alloy surface during wear results in work-hardening behavior, consistent with previous reports on the sliding wear of Ti alloys [39]. In contrast, the higher hardness of the Tens-20-A specimen leads to smaller grooves, less debris and reduced adhesion of particles to the surface. Figure 13i,j demonstrate the distribution of smaller particles on the surface of the wear patterns. This observation suggests the presence of an abrasive wear mechanism during the friction process of the Ti-15Mo alloy. Higher hardness results in worn tracks with shallower grooves. However, in a reciprocating sliding process, the surface undergoes repeated work hardening, leading to the development of layered cracks that eventually propagate and contribute to layered cracking of the material. The motion of the counter ball induces plastic deformation, generating abrasion particles that contribute to wear debris or form a transfer layer. Accumulated wear debris forms significant grooves on the surface, with a varying width of the wear tracks based on the alloy strength, exhibiting undulation and differences in width. The worn surfaces of all specimens display shallow grooves, confirming the presence of an abrasive wear mechanism [40].

4. Discussion

Different from traditional solution and aging treatment methods [41], based on the twinning-induced plasticity (TWIP) effect, {332} twin structures were introduced into the Ti-15Mo alloy through pre-tension deformation. Numerous HABs via twinning in the matrix have refined the initial grain formed during the ST. Based on these refined sub-structures and induced ω phases, the finer α phases at the nanoscale were precipitated during the aging treatment and distributed homogeneously. The microstructure of the alloy after pre-tension and aging enhanced the passivation effect and corrosion resistance of the alloy in the simulated human environment. The microstructural refinement via {332} twins and finer precipitates improved the hardness of the alloy and enhanced the wear resistance during friction in PBS solution. A schematic diagram (in Figure 14) has been drawn to depict the passivation behaviors of the ST-A and Tens-20-A alloy based on the relationship between the microstructure, passivation film and corrosion resistance. In PBS solution, the passivation film on the biomedical Ti-15Mo alloy mainly comprises the oxidation of Ti and Mo, initiated at surface flaws. In contrast to the limited element diffusion in the ST-A alloy (as Figure 14a), the extensive boundaries of the ultra-fine alloy offer numerous sites for the nucleation and growth of the passivation film (as Figure 14). The ultra-fine alloy is capable of forming a thicker film containing Ti and Mo oxide, enhancing the corrosion resistance of the alloy. Initially, the passivation film enables the creation of a more stable and efficient oxide layer to isolate electron and ion exchange, despite the existence of donor concentrations (including oxygen vacancy (O_vacancy) and metallic vacancy (M_vacancy)). Furthermore, a thicker oxide layer can prolong the corrosive and dissolution processes of the metal surface through the passivation film. Ultimately, the presence of a superior oxide film can impede the corrosion reaction, thereby enhancing the corrosion resistance of the ultra-fine alloy.

5. Conclusions

In this study, the effects of finer α phases within {332} twin structures via pre-tension and aging on the corrosion and wear resistance of a Ti-15Mo alloy were investigated systematically. The conclusions are as follows:

• 1.. The refined sub-structures of the Ti-15Mo alloy with {332} twins via pre-tension played the role of restricting the growth of α phases during aging, resulting in forming finer needle-shaped α precipitates. With the increase in pre-tension, finer {332} twins of the Ti-15Mo alloy after pre-deformation contributed to decreasing the size of the α precipitates formed during the aging process. The hardness of the aged alloy after pre-tension also increased from 272.0 HV in the ST-A specimen to 335.4 HV in the Tens-20-A specimen.

• 2.. The microstructure of the alloy after pre-tension and aging is finer and has more boundaries, which enhanced the passivation film structure and corrosion resistance of the alloy. The thickness of the passivation film on the surface of the aged alloy after pre-tension increases with the increase in the pre-deformation. The improved passivation film structure on the surface of the alloy after pre-deformation and aging enhanced the corrosion resistance. The I_pass value of the Tens-20-A specimen was 0.327 μA/cm², less than 1.382 μA/cm² of the ST-A specimen, increasing the corrosion resistance by 4.23 times.

• 3.. The worn track depth and debris size of the aged alloy after pre-tension decreased with the increase in pre-deformation. Compared with the ST-A specimen, which had a wear loss of 3.32 × 10⁷ μm³, the alloy after pre-tension and aging shows a better wear resistance. The wear loss of the Tens-20-A specimen in the same conditions was 2.24 × 10⁷ μm³, and the wear resistance was enhanced by 33%.

Author Contributions

Methodology—writing and original draft, S.Z.; conceptualization—review and editing, F.Y.; methodology, W.S.; investigation, N.C. and T.X. All authors have read and agreed to the published version of the manuscript.

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

Data are contained within the article.

Conflicts of Interest

The authors declare no conflicts of interest.

Footnotes

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Figures and Tables

Figure 1. The schematic diagram of the experimental process.

Figure 2. OIM and PFs of the Ti-15Mo alloy treated at 790 °C for 1 h. (a) OIM, (b) PFs. The scanning step size was 2 μm.

Figure 3. OIMs and BCD of the Ti-15Mo alloy after pre-tension with 3%, 9%, 15% and 20% strain. (a) 3%, (b) 6%, (c) 15%, (d) 20%. The scanning step size was 0.1 μm.

Figure 4. Grain size distribution of the Ti-15Mo alloy after different pre-tensions.

Figure 5. BF and DF TEM microstructure of the Ti-15Mo alloy after pre-tension with 9% and 20% strain. (a) BF with 9% pre-tension, (b) SAED with 9% pre-tension, (c) BF and SAED with 20% pre-tension.

Figure 6. KAM and {332} boundaries’ synthetic images of the Ti-15Mo alloy after pre-tension: (a) Tens-3, (b) Tens-9, (c) Tens-15 and (d) Tens-20 sample. The scanning step of EBSD was 0.1 μm.

Figure 7. (a) SEM and (b) TEM images of ST-A specimen.

Figure 8. SEM and TEM microstructure of the Ti-15Mo alloy after pre-tension and aging: (a,e) Tens-3-A specimen, (b,f) Tens-9-A specimen, (c,g) Tens-15-A specimen, (d,h) Tens-20-A specimen, (a–d) SEM microstructure, (e–h) TEM microstructure.

Figure 9. (a) OCP curves, (b) 3600 s potential of ST-A, Tens-3-A, Tens-9-A, Tens-15-A and Tens-20-A specimens in PBS solution.

Figure 10. (a) Nyquist plots and (b) Bode and impedance plots of ST-A, Tens-3-A, Tens-9-A, Tens-15-A and Tens-20-A specimens in PBS solution.

Figure 11. (a) M-S data and fitting curves and (b) potential dynamic polarization curves of ST-A, Tens-3-A, Tens-9-A, Tens-15-A and Tens-20-A specimens in PBS solution.

Figure 12. (a) Friction coefficient, (b) hardness and wear loss of Ti-15Mo alloy performed by different processes.

Figure 13. Three-dimensional morphology of the worn tracks on (a,b) ST-A specimen, (c,d) Tens-3-A specimen, (e) and SEM images of the worn tracks on (f) Tens-9-A specimen, (g,h) Tens-15-A specimen, (i,j) Tens-15-A specimen.

Figure 14. The schematic diagram of corrosion behavior and passive film of the Ti-15Mo alloy in PBS solution: (a) after solution treatment and aging, (b) after pre-deformation and aging.

Table 1

The fitted values of R_s, R_f, CPE, n and calculated values of C_eff, L_ss of Ti-15Mo alloy performed with different processes.

Specimen	Rs (Ω·cm2)	Rf (kΩ·cm2)	CPEf (μΩ−1·cm−2·sn)	n	Ceff (μF·cm−2)	LSS (nm)
ST-A	52.4	843	6.53	0.837	9.10	6.32
Tens-3-A	38.3	243	5.75	0.851	6.10	9.43
Tens-9-A	32.0	228	3.94	0.816	3.85	15.00
Tens-15-A	32.8	123	2.93	0.834	2.39	24.12
Tens-20-A	33.2	187	2.43	0.834	2.07	27.71

Table 2

The slope values of fitted M-S curves and calculated N_D values of passive film on the Ti-15Mo alloy after different pre-deformation and aging.

Specimen	b (1010)	ND (1020 cm−3)
ST-A	2.06	11.4
Tens-3-A	2.85	8.26
Tens-9-A	2.59	9.09
Tens-15-A	2.81	8.38
Tens-20-A	3.75	6.28

Table 3

The values of E_corr, I_corr and I_pass of the Ti-15Mo alloy performed by different processes in PBS solution.

Specimen	Ecorr(V)	Icorr(μA/cm2)	1st Ipass(μA/cm2)@0 mV	2nd Ipass(μA/cm2)@1000 mV
ST-A	−0.49	0.89	1.38	5.04
Tens-3-A	−0.40	0.15	0.42	4.99
Tens-9-A	−0.47	0.19	0.73	6.90
Tens-15-A	−0.51	0.28	0.80	6.55
Tens-20-A	−0.48	0.20	0.33	3.63

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