Tribocorrosion response of Fe-modified titanium alloys in simulated body fluids with varying glucose concentrations

Abstract

This study evaluates the tribocorrosion behavior of Fe-modified titanium alloys (Ti-3Fe, Ti-4.5Al-1V-3Fe, Ti-6Al-1V-3Fe) in comparison with Ti-6Al-4V in Hank’s balanced salt solution (HBSS) simulating normal (0 mg/dL), pre-diabetic (100 mg/dL) and diabetic (200 mg/dL) physiological glucose conditions. Tribocorrosion tests were conducted using a reciprocating sliding configuration under open-circuit potential (OCP) and potentiostatic control (+0.5 V vs Ag/AgCl) to assess passive film stability, depassivation–repassivation kinetics and wear–corrosion synergy. Mechanical sliding induced cathodic OCP transients due to passive film rupture, with Ti-3Fe showing the largest potential drop (0.61 V, from −0.42 to −1.03 V) and negligible recovery (<0.01 V), indicating poor repassivation. In contrast, Ti-6Al-1V-3Fe and Ti-6Al-4V exhibited anodic OCP recovery of +0.059 and +0.095 V, reflecting enhanced passive film reformation. Under potentiostatic conditions, Ti-3Fe recorded peak sliding current densities of 6.4 × 10⁻⁴ A, while Ti-6Al-1V-3Fe and Ti-6Al-4V remained below 4.3 × 10⁻⁴ A, indicating higher resistance to depassivation. Increasing glucose concentration progressively destabilized the electrochemical response, with OCP fluctuation ranges increasing from 0.05 V in normal HBSS to 0.09 V in diabetic HBSS, accompanied by elevated sliding current densities and reduced post-sliding recovery. Wear analysis showed the lowest specific wear rate for Ti-6Al-4V (1.68 × 10⁻⁶ mm³/N·m) in normal HBSS, while Ti-3Fe consistently exhibited the highest material loss across all environments. The results show that alloy composition and glucose concentration in physiological media significantly influence tribocorrosion behavior, with Al-V-containing alloys demonstrating superior electrochemical stability and wear resistance, while Fe-rich titanium alloys exhibit greater susceptibility, particularly under diabetic conditions.

Keywords

tribocorrosion titanium alloys passive film stability glucose-modified HBSS sliding-induced degradation

Introduction

Titanium and its alloys are extensively used in biomedical implants due to their high specific strength, suitable elastic modulus, excellent corrosion resistance and favorable biocompatibility, which collectively ensure reliable long-term performance in physiological environments.^1–4 Among these alloys, Ti-6Al-4V remains the most widely employed material for load-bearing orthopedic and dental applications, accounting for nearly 45% of global titanium alloy usage.⁴ Despite its widespread adoption, long-term clinical use of Ti-6Al-4V has raised concerns about the release of aluminum and vanadium ions, which have been associated with adverse biological responses, including inflammatory reactions and neurological complications.^4,5

To address these concerns, alternative titanium alloys have been developed by replacing conventional alloying elements with more biocompatible and economically viable β-stabilizers. Iron (Fe) has emerged as an attractive candidate due to its strong β-stabilizing capability, low cost and widespread availability compared with Nb, Ta and Mo.^6–8 Previous studies have demonstrated that controlled Fe additions can enhance solid-solution strengthening, refine microstructure and provide mechanical properties comparable to or exceeding those of Ti-6Al-4V, while maintaining acceptable corrosion resistance.^6,9–11 Bodunrin et al.⁶ developed a series of titanium alloys by systematically modifying the composition of the conventional Ti-6Al-4V alloy to reduce alloying costs while maintaining desirable mechanical and corrosion properties. Their findings showed that partial substitution of vanadium with iron, a low-cost and readily available β-stabilizing element, produced alloys with mechanical properties comparable to or exceeding those of Ti-6Al-4V.⁶ Furthermore, Fe-modified titanium alloys can be processed using conventional melting and thermomechanical routes, reducing manufacturing complexity and overall material cost. However, while their mechanical and corrosion properties have been reasonably documented, their degradation behavior under simultaneous mechanical and electrochemical loading remains insufficiently understood.

Recent research highlights the significance of tribocorrosion in biomedical titanium alloys and shows how physiological conditions influence degradation. Ti-6Al-4V demonstrates that hyperglycemic (high glucose) environments significantly increase material loss due to wear–corrosion synergism, suggesting that pathological conditions such as diabetes can worsen implant degradation compared with normal conditions.¹² Advanced analyses also reveal that the chemical composition of biological solutions affects tribocorrosive behavior of titanium alloys, indicating that electrolyte chemistry can alter passive film stability and synergistic material wear.^13,14 In addition, recent studies on novel alloy systems such as Ti–Nb–5Mo show that composition and microstructural design can significantly influence tribocorrosion resistance in simulated physiological environments.¹⁴ Finally, systematic reviews on surface treatment effects suggest that optimized surface modifications can improve tribocorrosion resistance by enhancing passive film robustness, which is critical for long-term implant durability.¹⁵

In biomedical applications, tribocorrosion studies are vital because orthopedic implants undergo simultaneous mechanical wear and electrochemical attack, creating a synergistic effect in which total material loss significantly exceeds the sum of the individual processes.^16,17 This research is justified because mechanical sliding ruptures the protective passive film on the metal surface; if repassivation is slow, corrosion is accelerated, leading to the release of metallic ions and wear debris that can trigger adverse biological reactions and shorten implant service life.^18,19 Within the investigated alloys, the roles of alloying elements are important but not absolute. Iron (Fe) serves as an attractive, low-cost β-stabilizer that enhances solid-solution strengthening and promotes microstructural refinement,^20,21 while its influence on passive film stability remains complex and context-dependent. Aluminum (Al) contributes to surface protection by promoting the formation of Al₂O₃-enriched mixed oxide layers that improve film adhesion and chemical stability.²² Vanadium (V) enhances repassivation kinetics and supports the formation of compact TiO₂-based passive films, enabling rapid recovery of the protective surface state after mechanical damage.²³

Prediabetic and diabetic environments are characterized by persistently elevated glucose levels, which can influence implant degradation beyond normal physiological conditions. High glucose concentrations can alter electrochemical reactions, compromise passive film stability and intensify wear–corrosion interactions on titanium and Fe-modified alloys, increasing corrosion susceptibility.^24,25 Glucose and reactive species can disrupt oxide film formation and retard repassivation, while chronic inflammation and oxidative stress in diabetic patients further impair healing and increase implant complications.²⁶ Given the rising prevalence of diabetes and its association with implant failure, understanding tribocorrosion under glucose-rich conditions is critical; however, systematic studies on the effect of glucose concentration on passive film stability and wear–corrosion synergy in Fe-modified titanium alloys remain limited.

Therefore, this study investigates the tribocorrosion behavior of Ti-3Fe, Ti-4.5Al-1V-3Fe, Ti-6Al-1V-3Fe and Ti-6Al-4V alloys in Hank’s Balanced Salt Solution under normal, pre-diabetic and diabetic glucose conditions. Tribocorrosion tests conducted under open-circuit and potentiostatic regimes are used to evaluate electrochemical stability, passive film recovery and wear mechanisms, to assess the suitability of Fe-modified titanium alloys for biomedical applications in glucose-rich physiological environments.

Experimental procedure

Alloy processing

The titanium alloys investigated in this study (Ti-3Fe, Ti-4.5Al-1V-3Fe, Ti-6Al-1V-3Fe and Ti-6Al-4V) were synthesized via vacuum induction melting furnace from high-purity elemental constituents (99.9 wt.%). The process was conducted on a water-cooled copper hearth within a chamber evacuated to 5 × 10⁻³ Pa and subsequently backfilled with high-purity argon (99%). Each alloy was turned and subjected to multiple remelting cycles to ensure chemical homogeneity. Following casting, all alloys were subjected to hot isostatic pressing (HIP) under an argon atmosphere, followed by furnace-cooling to room temperature to eliminate residual porosity. The alloys were then prepared into specimens suitable for subsequent tribocorrosion testing.

Sample preparation

Alloy specimens were precision-machined to a diameter of 20 mm and a thickness of 3 mm, then mounted in epoxy resin for tribocorrosion testing. The exposed surface was sequentially ground using silicon carbide (SiC) abrasive papers up to 2000 grit, followed by polishing with a 3 μm diamond paste and final finishing with a colloidal silica suspension to achieve a mirror-like metallographic surface. Before testing, the specimens were ultrasonically cleaned in deionized water to remove residual contaminants and air-dried under ambient laboratory conditions. Although the specimens had a diameter of 20 mm, only a centrally located circular region with a diameter of 6 mm was exposed during tribocorrosion experiments by covering the remaining surface with epoxy resin, resulting in an effective exposed electrode area of 0.283 cm².

Tribocorrosion testing setup

Tribocorrosion experiments were carried out using a ball-on-disk tribometer (CSM Instruments, Switzerland) integrated with a potentiostat/galvanostat/ZRA system (Gamry Reference 600+). The combined setup enabled potentiostatic control and simultaneous evaluation of mechanical degradation and electrochemical behavior. A conventional three-electrode configuration was employed, with the alloy specimen serving as the working electrode (WE), a platinum wire as the counter electrode (CE) and an Ag/AgCl (3 M KCl) electrode as the reference electrode (RE). Sliding tests were conducted using a 6 mm alumina (Al₂O₃) ball as the counterbody at a constant sliding speed of 0.08 m/s over a total sliding distance of 300 m, with an acquisition (wear track) radius of 5 mm and a sliding/data acquisition frequency of 2 Hz.

Testing conditions

A schematic tribocorrosion mechanism comparison was developed to illustrate the distinct electrochemical–mechanical interactions under open-circuit potential (OCP) and potentiostatic conditions (Figure 1). Experiments were conducted in Hank’s Balanced Salt Solution (HBSS), prepared to reflect physiological ionic composition, using modified HBSS containing calcium and magnesium (without phenol red) and anhydrous dextrose, both supplied by Sigma-Aldrich, USA, with solution pH maintained at 7.4 ± 0.1 and temperature at 37°C ± 1°C. Three glucose concentrations were used to simulate physiological states: normal (0 mg/dL), pre-diabetic (100 mg/dL) and diabetic (200 mg/dL). Alloy surfaces were electrochemically cleaned at −1 V for 15 min before testing. Experiments were performed under OCP and potentiostatic regimes, where OCP tests involved 75 min stabilization with alternating 15 min sliding/rest cycles to monitor spontaneous passive film rupture and recovery via potential evolution, while potentiostatic tests applied +0.5 V versus Ag/AgCl for 100 min (30 min pre-sliding stabilization, 30 min active sliding, 40 min recovery), with current transients used to quantify anodic film breakdown and reformation during mechanical wear. All tests were conducted under a normal load of 2 N, sliding speed of 0.08 m s⁻¹ and track radius of 5 mm, with coefficient of friction (COF), potential shifts and current fluctuations continuously recorded.

Figure 1.

Schematic comparison of alloy surface behavior under open-circuit potential (OCP) and potentiostatic (+0.5 V vs Ag/AgCl) conditions during tribocorrosion tests.

Data acquisition and analysis

During the tribological experiments, the open-circuit potential (OCP), current transients and coefficient of friction (COF) were recorded simultaneously to examine the interaction between mechanical wear and electrochemical corrosion processes. The total wear volume and specific wear rate (SWR) were determined through profilometric evaluation of the wear tracks. The chemical wear volume (Vchem) and mechanical wear volume (Vmech) were calculated from the integrated electrochemical charge transfer and profilometric displacement data, respectively.

Electrochemical data obtained from potentiodynamic polarization tests were further analyzed to assess chemical weathering behavior using the Vchem parameter. The calculation was based on the assumed stable ionic valence states of the constituent elements under physiological conditions (Ti⁴⁺, Al³⁺, V⁵⁺ and Fe²⁺/Fe³⁺), with corrosion potential and corrosion current density values used as inputs. The Vchem parameter was used as a relative indicator to compare the chemical weathering tendencies of the investigated alloys, rather than as an absolute measure of material degradation. To ensure reproducibility, each experimental condition was performed in duplicate.

Results and discussion

Tribocorrosion behavior of the developed alloys in normal HBSS

Figure 2(a)–(d) and Table 1 indicate that the alloys exhibit relatively stable open-circuit potentials (OCP) before sliding in normal HBSS (−0.42 to −0.51 V), corresponding to the formation of a passive oxide film on their surfaces in a simulated physiological environment.²⁷ The onset of sliding produces an immediate cathodic shift in OCP for all alloys due to mechanical rupture of the passive layer by the alumina counterbody, exposing fresh metal to the electrolyte and initiating depassivation. During sliding, the observed potential fluctuations reflect repeated depassivation–repassivation cycles, with their extent strongly dependent on alloy composition. Ti-3Fe (Figure 2(a)) shows the largest cathodic shift (0.61 V, from −0.42 to −1.03 V) and limited potential recovery, indicating severe passive film damage and sluggish repassivation, while Ti-4.5Al-1V-3Fe (Figure 2(b)) exhibits intermediate behavior characterized by intermittent and less stable potential oscillations.^28,29 In contrast, Ti-6Al-1V-3Fe (Figure 2(c)) and Ti-6Al-4V (Figure 2(d)) display more pronounced and sustained OCP fluctuations during sliding, suggesting a more dynamic passive film capable of rapid reformation under mechanical disruption. Notably, both Al–V-containing alloys recover to post-sliding OCP values comparable to or higher than their initial potentials (0.059 V for Ti-6Al-1V-3Fe and 0.095 V for Ti-6Al-4V), indicating enhanced passive film stability and improved resistance to tribocorrosion in HBSS compared with the Fe-rich compositions.^16,30,31

Figure 2.

Open circuit potential (OCP) measurements over time obtained during the tribocorrosion test in normal HBSS (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

Table 1.

Mean Open Circuit Potential (E_OCP) values for Ti-alloys measured in normal HBSS at three distinct phases: initial, sliding and final.

Alloys	OCP₀ (V)		OCP_sliding (V)		OCP_final (V)
Alloys	Stage 1	Stage 2	Stage 1	Stage 2	Stage 1	Stage 2
Ti-3Fe	−0.409	−0.582	−1.07	−1.085	−0.582	−0.625
Ti-4.5Al-1V-3Fe	−0.481	−0.494	−0.956	−0.774	−0.494	−0.517
Ti-6Al-1V-3Fe	−0.456	−0.365	−0.889	−0.914	−0.365	0.059
Ti-6Al-4V	−0.505	−0,587	−0.790	−0.709	−0.587	0.95

Evolution of anodic current (potentiostatic) during tribocorrosion

Potentiostatic tribocorrosion tests were conducted at +0.5 V to evaluate the stability of the passive film on the alloys under electrochemically controlled conditions. As shown in Figure 3, all alloys underwent cathodic reduction during the first 300 s, followed by a transient anodic response upon polarization. After a stabilization period of approximately 1200 s, sliding initiated a sharp increase in current, indicating depassivation of the surface oxide and exposure of the bare metal to the electrolyte.^31,32 Such behavior is consistent with recent reports that titanium alloys experience electrochemical instability under tribological loading, where passive film breakdown leads to localized oxidation and dissolution before repassivation occurs.³³

Figure 3.

Evolution of current as a function of testing time in normal HBSS: (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

During sliding, all alloys exhibited current fluctuations, reflecting repeated depassivation–repassivation cycles characteristic of tribocorrosion. However, full repassivation was not achieved due to continuous mechanical abrasion of the passive film.^31,34 Ti-3Fe exhibited more pronounced tribocorrosion transients compared with Ti-6Al-1V-3Fe and Ti-6Al-4V, signifying slower repassivation kinetics and greater susceptibility to anodic dissolution.³⁴ These findings align with recent studies demonstrating that alloying and microstructural modification have a significant influence on the electrochemical stability of titanium alloys under tribocorrosion conditions.³² The average current values before sliding (I₀), during sliding (I_sliding) and after sliding (I_final) are summarized in Table 2, highlighting the distinct electrochemical responses among the alloys.

Table 2.

Average values of the current measured before (I₀), during (I_sliding) and after sliding (I_final) of the titanium alloys sliding against an alumina ball in normal HBSS condition.

Alloys	I₀ (A)	I_sliding (A)	I_final (A)
Ti-3Fe	2.75E−06	6.40E−04	1.05E−06
Ti-4.5Al-1V-3Fe	2.29E−06	4.60E−04	7.13E−07
Ti-6Al-1V-3Fe	5.12E−06	4.22E−04	−6.72E−10
Ti-6Al-4V	1.78E−05	4.32E−04	−6.99E−09

Ti-3Fe exhibited the highest average current of 6.40 × 10⁻⁴ A during sliding, indicating an increased susceptibility to anodic dissolution. In contrast, Ti-4.5Al-1V-3Fe, Ti-6Al-1V-3Fe and Ti-6Al-4V displayed comparable average current values, with Ti-6Al-1V-3Fe showing the lowest at 4.22 × 10⁻⁴ A, which suggests a reduced oxidation and dissolution rate. After 3600 s, all alloys demonstrated a significant decline in current, attributed to anodic repassivation and the formation of a thicker passive film in less-worn regions. Notably, Ti-6Al-1V-3Fe and Ti-6Al-4V exhibited a more pronounced reduction in post-sliding current, indicating the development of a more stable and adherent passive film compared to Ti-3Fe and Ti-4.5Al-1V-3Fe.^35–37

Coefficient of friction during potentiostatic tribocorrosion test

Figure 4 shows that the alloys undergo an initial running-in phase before stabilizing into a steady-state friction regime, which occurs after approximately 1500 s. During this steady state, the coefficient of friction (COF) fluctuates due to the third-body effect and stick-slip behavior,³⁸ which are caused by the periodic accumulation and removal of wear debris at the sliding interface. Ti-3Fe, with the lowest average COF of 0.12, exhibits prominent sharp spikes. This behavior is directly attributable to its lower hardness compared to the other alloys, which leads to greater plastic deformation and abrasive wear, generating more wear debris that becomes entrapped and causes transient increases in friction.² In contrast, the Ti-4.5Al-1V-3Fe, Ti-6Al-1V-3Fe and Ti-6Al-4V alloys, which contain aluminum and vanadium, exhibit higher hardness and thus generate less debris. The alloys showed similar tribological behavior with average COF values of 0.26, 0.23 and 0.18. Their enhanced resistance to wear led to a more stable frictional response with fewer spikes, as the dynamics of wear particles fundamentally govern friction at the interface.^39,40

Figure 4.

Evolution of the coefficient of friction as a function of testing time under potentiostatic tribocorrosion conditions: (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

Wear volume loss under OCP and potentiostatic conditions in normal HBSS

The penetration depth profiles shown in Figure 5 and the tribocorrosion data in Table 3 illustrate the complex interactions among microstructural stability, passive film behavior and mechanical deformation in the titanium alloys studied under open circuit potential (OCP) in Hank’s Balanced Salt Solution (HBSS). The Ti-6Al-1V-3Fe alloy exhibited the deepest penetration depth (−8.35 × 10⁻⁵ mm) and the highest specific wear rate (2.53 × 10⁻⁶ mm³/N·m). This indicates an intensified tribocorrosion synergy due to continuous rupture and repassivation of the TiO₂ film during sliding. In contrast, the Ti-6Al-4V alloy exhibited a shallower penetration depth (−5.15 × 10⁻⁵ mm) and the lowest wear rate (1.68 × 10⁻⁶ mm³/N·m), reflecting superior tribocorrosion resistance associated with its well-balanced α + β microstructure, refined β-phase distribution and the ability of its compact TiO₂ passive layer to maintain electrochemical stability during sliding.^41,42 The intermediate responses of Ti-3Fe and Ti-4.5Al-1V-3Fe are attributable to the dual role of Fe: while it stabilizes the β-phase and enhances hardenability, it may also promote the formation of Fe-rich precipitates or β-islands that locally modify electrochemical activity and passive film stability, facilitating micro-corrosion and accelerating surface degradation under mechanical stress.⁴¹ These results indicate that changes in composition significantly impact the resilience of the passive film, deformation response and microstructural integrity, all of which affect the tribocorrosion behavior of titanium alloys in simulated physiological environments.

Figure 5.

Penetration depth versus wear track width profiles of the materials under OCP conditions in normal HBSS: (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

Table 3.

Tribocorrosion response parameters of the investigated alloys under open-circuit potential (OCP) conditions in normal HBSS.

Alloys	Worn surface area (mm²)	Max penetration depth (mm)	Total wear volume (mm³)	Specific wear rate (mm³/N.m)
Ti-3Fe	27.48	−7.63E−05	2.10E−03	3.49E−06
Ti-4.5Al-1V-3Fe	17.25	−7.70E−05	1.33E−03	2.21E−06
Ti-6Al-1V-3Fe	1824	−8.35E−05	1.52E−03	2.53E−06
Ti-6Al-4V	19.6	−5.15E−05	1.01E−03	1.68E−06

Figure 6 and Table 4 indicate that as aluminum (Al) and vanadium (V) are added, both the penetration depth and wear volume decrease, which demonstrates enhanced tribocorrosion resistance of the alloys in Hank’s Balanced Salt Solution (HBSS). The Ti-3Fe alloy exhibited the deepest penetration and the highest wear volume due to its less stable passive film and limited strengthening effect, which promotes active depassivation and accelerates the wear–corrosion interaction.³⁵ The incorporation of Al and V in the alloys Ti-4.5Al-1V-3Fe and Ti-6Al-1V-3Fe improved hardness and β-phase stability, resulting in a more homogeneous α + β structure and better adherence of the passive film. Al contributes to the formation of a protective Al₂O₃-enriched surface layer, while V enhances β-phase stabilization and repassivation kinetics, leading to reduced wear penetration and wear volume.⁴³ Furthermore, Ti-6Al-4V exhibited penetration profiles and specific wear rates like those of the Fe-modified alloys, suggesting that a moderate addition of iron (3 wt%) strengthens the β-phase and enhances tribocorrosion resistance without adversely affecting passivation behavior.⁴⁴ In addition, the wear decomposition data in Table 4 show that material loss under potentiostatic conditions is dominated by the chemical wear component (Vchem), with values of 6.35 × 10⁻² mm³ for Ti-3Fe and 4.19–4.56 × 10⁻² mm³ for the Al and V-containing alloys, while the mechanical component (Vmech) appears negative (−5.96 × 10⁻² mm³ for Ti-3Fe and −4.01 to −4.38 × 10⁻² mm³ for the other alloys), indicating that electrochemical dissolution is the controlling degradation mechanism and that mechanical sliding mainly promotes passive film breakdown and corrosion-assisted material loss.^45,46

Figure 6.

Penetration depth versus wear track width of the materials under potentiostatic conditions in normal HBSS: (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

Table 4.

Potentiostatic tribocorrosion behavior and electrochemical current response of the investigated alloys in normal HBSS.

Alloys	Area (mm²)	Pdmax (mm)	Total wear volume (mm²)	Vchem (mm³)	Vmech (mm³)	Specific wear rate (mm³/N.m)
Ti-3Fe	31.91	−1.22E−04	3.90E−03	6.35E−02	−5.96E−02	6.50E−06
Ti-4.5Al-1V-3Fe	21.69	−8.68E−05	1.88E−03	4.56E−02	−4.38E−02	3.14E−06
Ti-6Al-1V-3Fe	20.79	−8.71E−05	1.81E03	4.19E−02	−4.01E−02	3.02E−06
Ti-6Al-4V	21.83	−9.29E−04	2.03E−03	4.29E−02	−4.08E−02	3.38E−06

Wear track morphology in OCP and potentiostatic conditions in normal HBSS

In normal HBSS under OCP, SEM micrographs of the wear tracks (Figure 7) reveal that Ti-3Fe undergoes the most aggressive damage, exhibiting deep, wide grooves, debris accumulation and extensive delamination. This is typical of dominant abrasive/adhesive wear, where the passive film is repeatedly broken and slow to repassivate.^47,48 The Ti-4.5Al-1V-3Fe and Ti-6Al-1V-3Fe alloys exhibit narrower, more refined grooves and patches of residual oxide, consistent with enhanced surface hardness and more robust mixed films that reduce substrate exposure.⁴⁹ Among these, Ti-6Al-4V presents the mildest wear damage, with fine, shallow striations and minimal debris, indicating a stable and adherent oxide layer that limits metal transfer and suppresses tribocorrosion attack.³⁹ These findings corroborate earlier reports that Al and V improve the chemical stability and mechanical integrity of the passive film, thereby reducing the susceptibility of titanium alloys to tribocorrosion in simulated physiological solutions.⁵⁰

Figure 7.

SEM micrographs of the materials tested after OCP testing in normal HBSS: (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

The SEM images in Figure 8 reveal distinct tribocorrosion features across the studied alloys under potentiostatic conditions in normal HBSS. The Ti-3Fe alloy (Figure 8(a)) shows fine plowing lines and a relatively smooth wear track, indicating a wear mechanism influenced by passive film instability and corrosion-assisted material removal rather than purely mild abrasive wear. In comparison, the Ti-4.5Al-1V-3Fe and Ti-6Al-1V-3Fe alloys (Figure 8(b) and (c)) display broader grooves, micro-cracks and fragmented oxide layers, suggesting repeated rupture and reformation of the TiO₂ film due to electrochemical–mechanical interactions at the sliding interface. The deterioration becomes most severe in Ti-6Al-4V (Figure 8(d)), where deep grooves, oxide spallation and compacted debris signify an unstable tribolayer and a strong synergy between anodic dissolution and mechanical deformation. Such morphological evolution reflects the progressive destabilization of the TiO₂ passive film in alloys with increased α + β phase heterogeneity, which enhances local potential differences and accelerates corrosion wear synergy in chloride-containing electrolytes.^51,52 Similar findings have been reported by Du et al. and Fangaia et al.,^23,52,53 who demonstrated that galvanic interactions between α and β phases promote localized breakdown of the passive film during sliding in simulated body fluids. Likewise, Rundora et al.²³ reported that microstructural refinement leads to enhanced dielectric strength and improved adhesion of the passive film, resulting in reduced wear track depth and minimized material loss. On the other hand, potentiostatic polarization has been shown to promote oxide film rupture and intensify wear severity in Ti-6Al-4V alloys.⁵¹

Figure 8.

SEM micrographs of the materials tested after potentiostatic testing in normal HBSS: (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

Tribocorrosion behavior of titanium alloys in pre-diabetic HBSS

Figure 9 shows that Ti-3Fe experiences noticeable potential fluctuations during sliding, suggesting that glucose enhances repassivation. Ti-4.5Al-1V-3Fe, Ti-6Al-1V-3Fe and Ti-6Al-4V also show strong fluctuations, reflecting effective repassivation kinetics. Ti-6Al-1V-3Fe exhibits the highest pre-sliding potential, indicating superior passive-film stability. After sliding stops, all alloys spontaneously repassivate, but only Ti-3Fe recovers to a potential that exceeds its initial value, although slowly. These findings, supported by the quantitative data in Table 5, align with recent reports that glucose or biomolecules can stabilize passive films by facilitating dynamic repassivation in titanium alloys.^24,50

Figure 9.

Open circuit potential (OCP) measurements over time obtained during the tribocorrosion test in pre-diabetic HBSS: (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

Table 5.

Mean Open Circuit Potential (E_OCP) values for Ti-alloys measured in pre-diabetic HBSS at three distinct phases: initial, sliding and final.

Alloys	OCP₀ (V)		OCP_sliding (V)		OCP_final (V)
Alloys	Stage 1	Stage 2	Stage 1	Stage 2	Stage 1	Stage 2
Ti-3Fe	−0.532	−0.536	−1.08	−1.046	−0.536	−0.482
Ti-4.5Al-1V-3Fe	−0.418	−0.014	−0.921	−0.967	−0.014	−0.548
Ti-6Al-1V-3Fe	−0.154	−0.181	−0.458	−0.529	−0.181	0.177
Ti-6Al-4V	−0.435	−0,579	−1.059	−0.948	−0.579	0.613

Evolution of anodic current (potentiostatic) during tribocorrosion

Figure 10 and Table 6 show that, at the onset of sliding, all the alloys exhibited an abrupt increase in current, indicating the depassivation of the protective passive film. During the 3600 s sliding period, fluctuations in the current were observed, reflecting the continuous interplay between depassivation and repassivation processes. Isolated current peaks correspond to localized repassivation points that were rapidly broken down under sliding contact, a behavior commonly reported for titanium alloys in tribocorrosion studies.⁵⁴ As shown in Table 6, the Ti-3Fe alloy has the highest average current, indicating accelerated oxidation/dissolution processes and therefore a lower resistance to tribocorrosion.⁵⁵ Conversely, Ti-6Al-1V-3Fe attained the lowest average current, suggesting a superior ability to withstand tribocorrosion compared with the other alloys.⁵² Once sliding ceased, all alloys exhibited a current drop below their initial values, demonstrating spontaneous repassivation of their surfaces. Notably, Ti-6Al-1V-3Fe repassivated to the lowest residual current, further confirming its enhanced tribocorrosion resistance relative to the other tested alloys.⁵⁶

Figure 10.

Evolution of current as a function of testing time in pre-diabetic HBSS: (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

Table 6.

Average values of the current measured before (I₀), during (I_sliding) and after sliding (I_final) of the titanium alloys sliding against an alumina ball in pre-diabetic HBSS condition.

Alloys	I₀ (A)	I_sliding (A)	I_final (A)
Ti-3Fe	3.23−06	6.91E−04	5.01E−10
Ti-4.5Al-1V-3Fe	2.34E−06	5.00E−04	6.96E−07
Ti-6Al-1V-3Fe	1.06E−04	3.27E−04	4.11E−10
Ti-6Al-4V	3.16E−06	4.92E−04	7.53E−07

Coefficient of friction during potentiostatic tribocorrosion test in pre-diabetic HBSS

Figure 11 shows the evolution of the coefficient of friction (COF) under OCP conditions. A run-in period, reflecting the initial adaptation of sliding interfaces, was evident for Ti-4.5Al-1V-3Fe, Ti-6Al-1V-3Fe and Ti-6Al-4V, but less distinct for Ti-3Fe.² Sharp fluctuations in COF observed for Ti-3Fe, Ti-4.5Al-1V-3Fe and Ti-6Al-1V-3Fe are attributed to the third-body effect, where wear debris becomes entrapped and intermittently disrupts the contact, leading to transient rises in friction.⁵⁷ In contrast, Ti-4.5Al-1V-3Fe, Ti-6Al-1V-3Fe and Ti-6Al-4V rapidly attained steady-state conditions, indicating more effective surface accommodation and tribo-oxide film stabilization. In comparison, Ti-3Fe struggled to achieve stability, likely due to its lower alloying element content, which limited the formation of protective layers.⁵⁸ The average COF of all alloys remained below 0.45, consistent with metal–ceramic tribocouples with Ti-6Al-4V showing the highest steady value (0.40), suggesting higher susceptibility to debris accumulation and tribo-oxidation during sliding.⁵⁹

Figure 11.

Evolution of the coefficient of friction as a function of testing time under potentiostatic tribocorrosion conditions for pre-diabetic HBSS: (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

Wear volume loss under OCP and potentiostatic conditions in pre-diabetic HBSS

In pre-diabetic Hank’s Balanced Salt Solution (HBSS), the tribocorrosion resistance of the investigated alloys is significantly influenced by their specific elemental compositions, as illustrated by the wear track profiles in Figure 12(a)–(d). The Ti-3Fe alloy (Figure 12(a)) and the Ti-4.5Al-1V-3Fe alloy (Figure 12(b)) exhibited intermediate wear responses compared to the more complex formulations, while Ti-6Al-1V-3Fe (Figure 12(c)) demonstrated superior resistance characterized by the narrowest wear track and the lowest specific wear rate of 1.65 × 10⁻⁶ mm³/N.m. Conversely, the commercial Ti-6Al-4V alloy (Figure 12(d)) suffered the deepest penetration of −8.09 × 10⁻⁵ mm and the highest wear volume, as detailed in Table 7. These findings align with the tribocorrosion interaction model, which attributes total material loss to the dynamic balance between mechanical wear and electrochemical dissolution governed by the stability and repassivation ability of the oxide film.^14,18 While the presence of glucose in the solution can further destabilize this passive layer and increase electrochemical activity during sliding,⁶⁰ the enhanced performance of the Fe-modified alloys suggests that their compositional design promotes the formation of a denser and more chemically stable TiO₂-based layer, effectively limiting oxide rupture and localized metal dissolution.²¹ This improvement is consistent with previous reports regarding the increased compactness and adherence of passive films in simulated physiological environments.²³

Figure 12.

Penetration depth versus wear track width profiles of the materials under OCP conditions in pre-diabetic HBSS: (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

Table 7.

Tribocorrosion response parameters of the investigated alloys under open-circuit potential (OCP) conditions in pre-diabetic HBSS.

Alloys	Worn surface area (mm²)	Max penetration depth (mm)	Total wear volume (mm³)	Specific wear rate (mm³/N.m)
Ti-3Fe	23.41	−6.95E−05	1.63E−03	2.71E−06
Ti-4.5Al-1V-3Fe	17.74	−6.54E−05	1.16E−03	1.93E−06
Ti-6Al-1V-3Fe	16.51	−5.98E−05	9.87E−04	1.65E−06
Ti-6Al-4V	20.61	−8.09E−05	1.67E−03	2.78E−06

Figure 13 and Table 8 present the potentiostatic tribocorrosion behavior of Ti-3Fe, Ti-4.5Al-1V-3Fe, Ti-6Al-1V-3Fe and Ti-6Al-4V alloys in pre-diabetic HBSS, where anodic polarization intensified electrochemical activity at the sliding interface, promoting passive film breakdown and accelerated anodic dissolution. Ti-3Fe showed the highest degradation susceptibility, with penetration depth of −1.19 × 10⁻⁴ mm, wear volume of 3.28 × 10⁻³ mm³ and specific wear rate of 5.47 × 10⁻⁶ mm³/N·m, indicating poor passive film stability, while the Al-V-Fe alloys exhibited significantly reduced wear volumes and rates due to superior protective oxide layer maintenance during sliding and commercial Ti-6Al-4V displayed intermediate performance. Wear decomposition data confirmed electrochemical (chemical) processes dominated degradation (high Vchem), with negative mechanical wear (Vmech) highlighting corrosion-driven removal facilitated by mechanical passive film destabilization rather than direct abrasion. The improved tribocorrosion resistance of Al-V-Fe modified alloys arises from alloying-induced microstructural refinement and enhanced, rapidly repassivating oxide film stability, minimizing tribocorrosion synergy. These results align with prior studies showing that potentiostatic polarization accelerates corrosion-assisted wear, while Al and V additions enhance tribocorrosion resistance in titanium alloys via more stable and adherent passive films.^19,61

Figure 13.

Penetration depth versus wear track width profiles of the materials under potentiostatic conditions in pre-diabetic HBSS: (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

Table 8.

Potentiostatic tribocorrosion behavior and electrochemical current response of the investigated alloys in pre-diabetic HBSS.

Alloys	Area (mm²)	Pdmax (mm)	Total wear volume (mm²)	Vchem (mm³)	Vmech (mm³)	Specific wear rate (mm³/N.m)
Ti-3Fe	27.53	−1.19E−04	3.28E−03	3.35E−01	−3.32E−01	5.47E−06
Ti-4.5Al-1V-3Fe	17.45	−7.97E−05	1.39E−03	4.58E−01	−4.57E−01	2.32E−06
Ti-6Al-1V-3Fe	20.64	−7.23E−05	1.49E03	2.81E−01	−2.80E−01	2.49E−06
Ti-6Al-4V	20.67	−1.01E−04	2.09E−04	4.42E−01	−4.40E−01	3.48E−06

Wear track morphology in OCP and potentiostatic conditions in pre-diabetic HBSS

The SEM images in Figure 14 illustrate the tribocorrosion features across the studied alloys under OCP conditions in pre-diabetic HBSS. The presence of elevated glucose in the HBSS simulates a physiologically challenging environment, potentially accelerating the corrosion rate.³³ The Ti-3Fe alloy (Figure 14(a)) exhibits the most severe degradation, characterized by a wide track with deep plow grooves and substantial wear debris, indicating aggressive material removal via abrasive wear.⁶² In contrast, Figure 14(b) and (c) exhibit significantly narrower and shallower wear tracks than those of Ti-3Fe, suggesting superior performance. This difference in damage is governed by the synergistic action of wear and corrosion, where mechanical sliding continuously removes the alloy’s protective oxide film (depassivation), exposing the reactive metal to the corrosive fluid.⁶³ The Ti-4.5Al-1V-3Fe and Ti-6Al-1V-3Fe alloys perform better because their compositional additions of Al and V promote more favorable repassivation kinetics or form a more stable surface layer in the HBSS, allowing them to better resist the combined mechanical and electrochemical attack.⁶⁴ This observation aligns with recent research indicating that simpler Ti-Fe compositions often have lower wear resistance compared to more advanced multi-component titanium alloys in simulated physiological environments.⁶²

Figure 14.

SEM micrographs of the materials tested after OCP testing in pre-diabetic HBSS: (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

The SEM micrographs in Figure 15 reveal distinct tribocorrosion morphologies across the tested alloys under potentiostatic conditions in pre-diabetic HBSS. The Ti-3Fe alloy (Figure 15(a)) exhibits a relatively smooth wear track with shallow grooves, indicating stable passive film formation and limited oxide rupture. As alloying additions increase, the Ti-4.5Al-1V-3Fe and Ti-6Al-1V-3Fe samples (Figure 15(b) and (c)) show moderately widened grooves and patches of delamination, suggesting increased mechanical wear coupled with localized depassivation.⁶⁵ In contrast, the Ti-6Al-4V alloy (Figure 15(d)) displays severe surface degradation with evidence of adhesive smearing and oxide spallation, attributed to enhanced electrochemical activity under anodic polarization.⁶⁶

Figure 15.

SEM micrographs of the materials tested after potentiostatic testing in pre-diabetic HBSS: (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

Tribocorrosion behavior of titanium alloys in diabetic HBSS

In diabetic HBSS (200 mg/dL glucose), the evolution of the open-circuit potential (OCP) reflects the superposition of rapid mechanically induced depassivation–repassivation events and slower chemically driven interfacial processes associated with glucose-rich environments. As shown in Figure 16(a)–(d) for Ti-3Fe, Ti-4.5Al-1V-3Fe, Ti-6Al-1V-3Fe and Ti-6Al-4V, respectively, all alloys exhibit an abrupt cathodic OCP shift immediately upon sliding initiation at 900 s, which occurs over a short time scale and is attributed to instantaneous mechanical rupture of the passive oxide film and exposure of fresh metal to the electrolyte, a characteristic response of tribocorrosion systems under sliding contact.¹⁸ The magnitude of these fast mechanical transients is composition-dependent, as summarized in Table 9, with the Fe-rich Ti-3Fe alloy (Figure 16(a)) showing the largest cathodic drop, indicating reduced resistance of its passive film to mechanical disruption, whereas the Al- and V-containing alloys (Figure 16(b)–(d)) exhibit relatively moderated transients consistent with improved film stability and repassivation kinetics. Superimposed on these fast sliding-induced OCP changes is a slower baseline evolution observed during prolonged exposure in diabetic HBSS, evident from the gradual pre-sliding stabilization and incomplete post-sliding recovery in Figure 16(a)–(d). This slow drift is attributed to the adsorption of glucose and glucose-derived organic species on the oxide surface, which progressively modifies passive film chemistry and shifts the corrosion potential independently of instantaneous mechanical events, as reported for hyperglycemic physiological environments.^14,23

Figure 16.

Open circuit potential (OCP) measurements over time obtained during the tribocorrosion test in diabetic HBSS: (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

Table 9.

Mean Open Circuit Potential (E_OCP) values for Ti-alloys measured in diabetic HBSS at three distinct phases: initial, sliding and final.

Alloys	OCP₀ (V)		OCP_sliding (V)		OCP_final (V)
Alloys	Stage 1	Stage 2	Stage 1	Stage 2	Stage 1	Stage 2
Ti-3Fe	−0.319	−0.474	−1.038	−0.967	−0.474	−0.486
Ti-4.5Al-1V-3Fe	−0.341	−0.548	−1.047	−1.042	−0.548	−0.559
Ti-6Al-1V-3Fe	−0.218	−0.605	−1.084	−1.098	−0.605	−0.679
Ti-6Al-4V	−0.72	−0.76	−1.14	−1.18	−0.76	−0.77

Evolution of anodic current (potentiostatic) during tribocorrosion

Figure 17 illustrates that the initiation of sliding at 900 s resulted in a sudden anodic (more noble) current shift. This shift is consistent with the mechanical removal of the passive oxide layer, exposing the active metal at the contact point. During the subsequent sliding interval of 3600 s, the current showed near-steady but fluctuating values, which indicates repeated depassivation and repassivation events occurring at the worn surface.³⁹ As detailed in Table 10, the higher average currents recorded during sliding are associated with an increased oxidation and dissolution (chemical) contribution to material loss under tribocorrosion conditions.⁶⁷ Although the alloys demonstrated broadly similar average currents in diabetic Hank’s Balanced Salt Solution (HBSS), Ti-3Fe exhibited the highest average current at 0.0007 A, while Ti-6Al-1V-3Fe showed the lowest. Notably, the latter alloy also stabilized at the most negative final current, which suggests a more stable and adherent passive film, indicating superior tribocorrosion resistance.⁵¹ These trends in alloy performance align with existing literature, which indicates that glucose and other organic compounds present in simulated biological fluids can significantly affect passive film stability and enhance electrochemical activity.⁶⁸

Figure 17.

Evolution of current as a function of testing time in diabetic HBSS: (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

Table 10.

Average values of the current measured before (I₀), during (I_sliding) and after sliding (I_final) of the titanium alloys sliding against an alumina ball in diabetic HBSS condition.

Alloys	I₀ (A)	I_sliding (A)	I_final (A)
Ti-3Fe	3.08E−06	0.0007	9.78E−07
Ti-4.5Al-1V-3Fe	2.29E−06	0.0006	5.7E−07
Ti-6Al-1V-3Fe	2.71E−05	0.0004	−3.85E−09
Ti-6Al-4V	1.70E−05	0.0005	6.28E−06

Coefficient of friction during potentiostatic tribocorrosion test in diabetic hbss

The evolution of the coefficient of friction (µ) during potentiostatic tribocorrosion tests in simulated diabetic Hank’s Balanced Salt Solution (HBSS), as illustrated in Figure 18, shows distinct tribological behavior based on alloy composition. Ti-6Al-4V exhibits the highest and most volatile COF, indicating that its passive layer is highly susceptible to mechanical and electrochemical degradation, leading to poor wear resistance.⁶⁹ In sharp contrast, the experimental β-stabilized alloy, Ti-4.5Al-1V-3Fe, demonstrates superior tribological stability, exhibiting the lowest and most stable COF, suggesting that the inclusion of iron (Fe) and the resulting microstructure promote the formation of a durable, low-shear-strength tribo-layer that minimizes material loss.⁶² The alloy Ti-6Al-1V-3Fe also performs well, maintaining a stable, intermediate COF. Finally, the binary alloy Ti-3Fe provides insight into the stabilization mechanism, starting with a relatively high, fluctuating COF that then slightly decreases and stabilizes, a characteristic response during the “running-in” period as initial surface asperities are smoothed and a partially protective film is established.⁴²

Figure 18.

Evolution of the coefficient of friction as a function of testing time under potentiostatic tribocorrosion conditions for diabetic HBSS.

Wear volume loss under OCP and potentiostatic conditions in diabetic HBSS

Under OCP conditions in diabetic HBSS, the penetration depth wear track width profiles shown in Figure 19(a)–(d) and the tribocorrosion parameters in Table 11 reveal a clear compositional influence on the degradation behavior of the alloys. The profile of Ti-3Fe (Figure 19(a)) exhibits a broader and deeper wear track, corresponding to the highest wear volume (2.19 × 10⁻³ mm³) and specific wear rate (3.65 × 10⁻⁶ mm³/N·m), indicating greater surface deterioration. Ti-4.5Al-1V-3Fe (Figure 19(b)) and Ti-6Al-1V-3Fe (Figure 19(c)) show narrower and shallower profiles, while Ti-6Al-4V (Figure 19(d)) presents the most stable track geometry, consistent with the lowest wear volume (1.28 × 10⁻³ mm³) and specific wear rate (2.14 × 10⁻⁶ mm³/N·m). These variations suggest that the presence of Al and V promotes the formation of a dense, adherent TiO₂-based passive layer capable of sustaining repassivation under sliding conditions, thereby minimizing the wear–corrosion synergy. In contrast, Ti-3Fe exhibits limited passive film stability, resulting in repeated oxide rupture and enhanced metal dissolution. The diabetic HBSS medium, enriched with glucose and reactive oxygen species, alters interfacial electrochemistry and impairs oxide regeneration, which intensifies the degradation of alloys with weaker passivity.²³ The moderate wear response of Ti-4.5Al-1V-3Fe and Ti-6Al-1V-3Fe, as shown in Figure 19 and Table 11, confirms improved passivity and mechanical stability under glucose-rich physiological conditions.⁷⁰

Figure 19.

Penetration depth versus wear track width profiles of the materials under OCP conditions in diabetic HBSS: (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

Table 11.

Tribocorrosion response parameters of the investigated alloys under open-circuit potential (OCP) conditions in diabetic HBSS.

Alloys	Worn surface area (mm²)	Max penetration depth (mm)	Total wear volume (mm³)	Specific wear rate (mm³/N.m)
Ti-3Fe	24.74	−8.85E−05	2.19E−03	3.65E−06
Ti-4.5Al-1V-3Fe	18.72	−7.20E−05	1.35E−03	2.25E−06
Ti-6Al-1V-3Fe	19.74	−8.16E−05	1.61E−03	2.68E−06
Ti-6Al-4V	18.08	−7.10E−05	1.28E−03	2.14E−06

The potentiostatic tribocorrosion performance of the alloys in aggressive diabetic HBSS is governed by alloying elements that influence passive film dynamics.¹² Ti-3Fe showed the poorest resistance, with the highest specific wear rate (6.22 × 10⁻⁶ mm³/N·m) and maximum penetration depth (1.34 × 10⁻⁴ mm), indicating weak passive layer integrity and slow repassivation after mechanical disruption, leading to dominant wear-accelerated corrosion as seen in the severe wear track (Figure 20(a)). In contrast, Ti-4.5Al-1V-3Fe, Ti-6Al-1V-3Fe and Ti-6Al-4V, incorporating Al and V, exhibited superior stability with specific wear rates around 3.0 × 10⁻⁶ mm³/N·m (Table 12), due to rapid reformation of a protective film that suppresses electrochemical transients and minimizes the chemical wear component (Vchem), confining total degradation mainly to mechanical removal (Vmech), as evidenced by contained wear tracks (Figure 20(b)–(d)). Wear decomposition results (Table 12) confirm that degradation in diabetic HBSS is dominated by chemical wear (Vchem), driven by passive film destabilization in the glucose-rich and protein-containing medium, promoting anodic dissolution under potentiostatic conditions, while negative Vmech values indicate that corrosion-controlled processes prevail over direct mechanical abrasion.⁴²

Figure 20.

Penetration depth versus wear track width of the materials under potentiostatic conditions in diabetic HBSS: (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

Table 12.

Potentiostatic tribocorrosion behavior and electrochemical current response of the investigated alloys in diabetic HBSS.

Alloys	Area (mm²)	Pdmax (mm)	Total wear volume (mm²)	Vchem (mm³)	Vmech (mm³)	Specific wear rate (mm³/N.m)
Ti-3Fe	27.84	−1.34E−04	3.73E−03	3.56E−01	−3.52E−01	6.22E−06
Ti-4.5Al-1V-3Fe	20.91	−8.54E−05	1.79E−03	5.25E−01	−5.23E−01	2.98E−06
Ti-6Al-1V-3Fe	21.4	−8.87E−05	1.90E−03	3.49E−01	−3.47E−01	3.16E−06
Ti-6Al-4V	19.17	−9.36E−05	1.79E−03	4.39E−01	−4.37E−01	2.99E−06

Wear tracks morphology in OCP and potentiostatic conditions in diabetic hbss

Figure 21 presents the surface morphologies of the wear tracks for Ti-3Fe, Ti-4.5Al-1V-3Fe, Ti-6Al-1V-3Fe and Ti-6Al-4V alloys after tribocorrosion testing under open-circuit potential (OCP) conditions in diabetic HBSS. Distinct variations are observed across the alloys, reflecting the combined influence of microstructural stability and the altered electrochemical environment induced by elevated glucose concentration. The Ti-3Fe alloy (Figure 21(a)) exhibits relatively shallow grooves and a smooth track, indicating corrosion-assisted wear and limited oxidative degradation, which suggests that the Fe addition promotes a more protective passive film, even in an aggressive medium. In contrast, the Ti-4.5Al-1V-3Fe surface (Figure 21(b)) shows patches of delamination and fine debris, implying intermittent oxide rupture due to the mechanical removal of the passive layer. The Ti-6Al-1V-3Fe alloy (Figure 21(c)) reveals wider grooves with micro-plowing features, signifying a higher wear rate that may stem from unstable film regeneration in the glucose-enriched solution, while the Ti-6Al-4V surface (Figure 21(d)) displays severe deformation and adhesive smearing, characteristic of mixed adhesive–abrasive wear under compromised passivity. The presence of glucose molecules and oxidative species in diabetic HBSS can modify the ionic balance and hinder the formation of a compact TiO₂ layer, accelerating localized wear and metal ion dissolution.¹² A similar observation was reported by Liu et al.⁶⁸ found that glucose-containing electrolytes reduce the corrosion resistance of titanium alloys by destabilizing the passive film and increasing frictional contact damage.

Figure 21.

SEM micrographs of the materials tested after OCP testing in diabetic HBSS: (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

Figure 22 shows the SEM micrographs of the wear scars on Ti-3Fe, Ti-4.5Al-1V-3Fe, Ti-6Al-1V-3Fe and Ti-6Al-4V alloys after tribocorrosion testing under potentiostatic conditions in diabetic HBSS. The Ti-3Fe alloy (Figure 22(a)) reveals deep grooves, oxide delamination and micro-pits, indicating intensified material removal due to enhanced electrochemical dissolution under constant anodic potential. The Ti-4.5Al-1V-3Fe alloy (Figure 22(b)) displays relatively smooth tracks with shallower grooves and fewer damaged regions, implying that Al addition contributes to improved oxide stability and localized repassivation even when polarized. The Ti-6Al-1V-3Fe alloy (Figure 22(c)) exhibits a combination of abrasive wear marks and microcraters, indicating intermittent disruption of the passive film and partial tribochemical resistance. Among the studied alloys, Ti-6Al-4V (Figure 22(d)) exhibits narrow grooves and minimal oxide spallation, suggesting a more compact and adherent passive film that sustains less damage under electrochemical load. These findings show that the applied potential significantly influences wear morphology in diabetic HBSS by accelerating passive film breakdown through the interaction of glucose-induced reactive oxygen species and anodic dissolution, which together intensify the wear–corrosion synergy.¹² The sustained polarization, therefore, promotes deeper material loss and increased surface reactivity compared to conditions of free corrosion. Similar observations have been reported where potentiostatic control in glucose-rich simulated body fluids enhances oxide film instability and aggravates tribocorrosion degradation of titanium-based alloys.⁴²

Figure 22.

SEM micrographs of the materials tested after potentiostatic testing in diabetic HBSS: (a) Ti-3Fe, (b) Ti-4.5Al-1V-3Fe, (c) Ti-6Al-1V-3Fe, and (d) Ti-6Al-4V.

Mechanistic evolution in tribocorrosion across normal, pre-diabetic and diabetic physiological environments: Influence of alloy chemistry

The tribocorrosion behavior of the investigated titanium alloys (Ti-3Fe, Ti-4.5Al-1V-3Fe, Ti-6Al-1V-3Fe and Ti-6Al-4V) across normal, pre-diabetic and diabetic HBSS environments is governed by the coupled interaction between mechanically induced depassivation and environment-dependent electrochemical repassivation, with alloy chemistry and electrolyte composition jointly controlling degradation mechanisms. In all environments, sliding disrupts the native TiO₂-based passive film, exposing fresh metal to the electrolyte and inducing transient electrochemical activity; however, the extent of depassivation and the efficiency of subsequent repassivation vary systematically with glucose concentration and alloying additions. In normal HBSS, relatively gentle ionic conditions promote rapid oxide reformation, leading to dynamic but stable depassivation–repassivation cycles, consistent with classical tribocorrosion behavior of titanium alloys in physiological solutions.^71,72

As the environment transitions to pre-diabetic HBSS, increased glucose concentration introduces competitive adsorption phenomena at the metal–electrolyte interface, which modify surface charge transfer reactions and hinder passive film regeneration. Recent electrochemical studies have shown that glucose adsorption alters oxide growth kinetics and increases passive film defect density, resulting in prolonged depassivated states and enhanced electrochemical instability during sliding.²⁵ Under these conditions, alloys containing Al and V exhibit comparatively improved repassivation behavior due to the formation of compact mixed oxides that resist dissolution, whereas Fe-rich alloys experience delayed potential recovery associated with the formation of less stable iron-containing oxides and hydroxides.

In diabetic HBSS, where glucose concentration is highest, tribocorrosion mechanisms evolve toward sustained depassivation and intensified synergistic wear–corrosion interactions. Hyperglycemic environments have been reported to significantly increase tribocorrosive material loss and friction coefficients for titanium alloys by destabilizing passive films and extending the duration of bare-metal exposure during sliding.¹⁴ The elevated glucose content enhances electrochemical dissolution and disrupts oxide re-formation, particularly in Fe-containing alloys, while Al- and V-modified systems retain comparatively better electrochemical resilience despite the overall increase in degradation severity. Similar observations across different titanium-based biomaterials confirm that changes in solution chemistry associated with metabolic conditions play a decisive role in dictating tribocorrosion response beyond mechanical effects alone.⁷²

Overall, the comparative analysis across normal, pre-diabetic and diabetic environments demonstrates a progressive shift from mechanically dominated tribocorrosion in normal HBSS to chemically assisted and electrochemically sustained degradation in diabetic HBSS. The results highlight that alloy-dependent passive film chemistry strongly influenced by Al, V and Fe additions controls repassivation kinetics and resistance to synergistic wear-corrosion processes under disease-relevant physiological conditions.

Conclusions

This study provides a comprehensive evaluation of the synergistic effects of alloy composition and glucose concentrations on the tribocorrosion behavior of iron-modified titanium alloys. The investigation elucidates the critical role of alloying elements and the chemical environment in determining the integrity of the passive oxide layer during mechanical wear. The results demonstrate that the tribocorrosive performance of these materials is highly sensitive to the glucose levels present within the electrolyte environment. The major findings are summarized as follows:

The research demonstrates that while mechanical sliding induces depassivation across all investigated alloys, the chemical composition is the primary determinant of surface recovery. Alloys containing aluminum and vanadium (Ti-6Al-4V and Ti-6Al-1V-3Fe) exhibit superior repassivation kinetics and electrochemical stability. In contrast, the Ti-3Fe alloy exhibits substantial cathodic potential shifts and a limited ability to restore its protective TiO₂ layer following mechanical rupture.

Increasing glucose concentrations in the simulated body fluids progressively degrade the tribocorrosion resistance of the alloys. This degradation is driven by the competitive adsorption of glucose molecules at the sliding interface, which retards the reformation of the passive film and facilitates higher current densities, thereby intensifying anodic dissolution.

A definitive transition in degradation modes was identified as the glucose concentration in the HBSS environment increased. In normal HBSS, material loss is primarily governed by mechanical wear; however, in simulated diabetic conditions, the process evolves into a chemically assisted and electrochemically sustained degradation where the synergy between wear and corrosion accounts for most of the total material volume loss.

Among the tested materials, Ti-6Al-4V and the Fe-modified Ti-6Al-1V-3Fe alloy consistently demonstrated the highest resistance to wear-accelerated corrosion across all simulated environments. Conversely, the Ti-3Fe alloy exhibited the greatest susceptibility to material loss, particularly in high-glucose electrolytes, suggesting that a balanced α + β microstructure and specific alloying stabilizers are essential for maintaining tribocorrosion resistance in aggressive simulated physiological conditions.

Footnotes

ORCID iDs

Oluwasegun Falodun

Michael Bodunrin

Ethical considerations

The authors declare that no ethical approval was necessary for this study, as it did not make use of human participants, data, or tissues, nor did it involve the use of live animals.

Funding

The authors received no financial support for the research, authorship, and/or publication of this article.

Declaration of conflicting interests

The authors declared no potential conflicts of interest with respect to the research, authorship, and/or publication of this article.

Data availability statement

The datasets generated and analysed during the current study are available from the corresponding author upon reasonable request.

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