Influence of solutioning and ageing heat treatments on microstructure, mechanical properties, and tribological behavior of Ti-6Al-5Zr-0.5Mo-0.2Si alloy

ABSTRACT

The present study investigates the effect of heat treatment on the microstructure, mechanical, and wear characteristics of the Ti-6Al-5Zr-0.5Mo-0.2Si alloy. Solution heat treatment combined with aging near the β-transus temperature substantially modified the microstructure, resulting in notable enhancements in the mechanical properties. Solution annealing followed by ageing at 1050°C was highly effective, as it formed secondary α phases near the retained β phase. This refined microstructure contributed to a notable increase in tensile strength, achieving a higher strength of 1270 ± 10 MPa (29% more) and a microhardness of 360 ± 5 HV (14% more) than the as-received base material through transformation strengthening mechanisms. Additionally, the heat-treated alloys demonstrated superior wear resistance compared to their as-received condition. High-temperature wear testing at 200 °C showed the formation of a stable tribo-oxide layer on the alloy surface, which acted as a protective barrier against direct metal-to-metal contact. This layer significantly enhanced wear resistance, reaching 4.39 × 10⁻³ mm³/Nm, an improvement of 33% over the as-received base metal.

Keywords:
Tensile properties; Microstructure; Near α-Ti alloys; Wear behaviour

1. INTRODUCTION

Titanium (Ti) and its alloys are extensively utilized in aerospace applications due to their exceptional mechanical characteristics and corrosion resistance, particularly in high-temperature environments. Titanium alloys constitute approximately 1/3rd of the weight in modern aircraft engine construction, ranking as the second most prevalent material after nickel-based super alloys [¹]. The strength of titanium alloys is influenced by several factors, including alloy composition, microstructure, processing techniques, and heat treatment, which affects the phase stability [²]. Ti alloys are generally classified by their crystal structure: α Ti alloys possess a hexagonal close-packed (hcp) configuration, β Ti alloys exhibit a body-centred cubic (bcc) arrangement, and α + β Ti alloys consist of a mix of both phases. The α + β Ti alloys are further subdivided into near-α, α + β, and near-β types. Of these, α and near-α Ti alloys are the most frequently utilized due to their superior formability and toughness, especially at high temperatures [³,⁴,⁵,⁶].

Near-α titanium alloys, developed for high-temperature environments, are formulated using a Ti-Al-Zr-Mo-Si alloying structure, the concentration of β stabilizing elements (Mo and Si) is maintained below 2%, and the silicon content is limited to a maximum of 0.5%. These stabilizers result in a minor presence of β phase and silicide precipitates within the microstructure. Such alloys exhibit outstanding performance in high-temperature environments, especially in aerospace engines, owing to their excellent thermo-mechanical characteristics and creep resistance [⁷,⁸,⁹]. Due to their remarkable mechanical properties and thermal stability, numerous near-α titanium alloys have been developed and utilized in high-temperature settings, especially for aircraft engine parts. IMI 685 (Ti-6Al-5Zr-0.5Mo-0.2Si) is a key near-α titanium alloy engineered for high-temperature usage up to 823K in jet engines. In 1964, the IMI 685 alloy was incorporated into the Rolls-Royce Adour engine, which powered the Jaguar and Hawk aircraft, thanks to its superior creep resistance, high fracture toughness, and favourable weldability [¹⁰]. This material can effectively replace the compressor blades in the gas turbine application on the high-pressure compressor stages. ELSHAER et al. [¹¹] reported that strain hardening combined with aging treatments significantly affects the microstructural evolution and mechanical performance of cast titanium alloys. The alloy’s microstructure comprises primary α-phase (α), secondary α-phase (α″), and retained β-phase (β). After undergoing solution treatment, either by air cooling or water quenching, followed by an aging process, the secondary α-phase (α″) was found to form within the retained β-phase (β). The highest wear rate was detected in the water-quenched (WQ) specimen. In contrast, the lowest wear rate was noted for the specimen that underwent water quenching followed by aging (WQ + Aging), attributed to the increased hardness. Aging after solution treatment significantly improves wear resistance, and in the case of WQ + Aging compared to WQ alone, the enhancement in wear performance can be as much as approximately 125%. WANG et al. [¹²] and YANG et al. [¹³] reported that the solution treatment and aging (STA) process notably impacts the microstructure and mechanical properties of the Ti-6Al-4V alloy. The findings indicate that tailored heat treatment parameters can improve properties like strength and hardness by regulating the morphology and arrangement of the α and β phases. The findings offer a valuable guideline for tailoring heat treatment processes to achieve specific performance requirements in aerospace and other high-performance applications involving α+β titanium alloys.

LIU et al. [¹⁴] reported that short-duration duplex heat treatment has proven effective in enhancing the strength of titanium alloys. The process, involving high-temperature treatment around 1050°C followed by aging, helps the formation of a fine secondary α phase, which contributes significantly to the enhancement in mechanical strength. The study highlights that such heat treatment approaches can achieve enhanced strength without prolonged thermal exposure, making them suitable for industrial applications that necessitate both efficiency and performance. ZHANG et al. [¹⁵] reported that a bimodal microstructure characterizes the Ti-Al-Sn-Zr-Mo-Nb-W-Si alloy, which exhibits advantageous tensile properties at 700°C. Both primary alpha and transformed beta phases contribute to a balanced combination of strength and ductility at elevated temperatures. Their findings highlight the importance of microstructural design through appropriate heat treatment to optimize the elevated temperature performance of near-α titanium alloys, making the material suitable for demanding aerospace applications. ALI et al. [¹⁶] reported that the heat treatment significantly affects the mechanical performance and microstructural features of the near-α titanium alloy Ti60. Different heat treatment conditions altered phase distribution, notably affecting the morphology and organization of the α and β phases, directly influencing mechanical properties like strength, ductility, and creep resistance. Their work shows that suitable heat treatment can fine-tune these properties, making Ti60 well-suited for use in high-temperature aerospace environments.

WANG et al. [¹⁷] reported that the heat treatment has a pronounced effect on the microstructure and mechanical behaviour of selectively laser-melted Ti-6Al-4V (Ti64), influencing phase distribution, grain morphology, and overall performance. The study verified that appropriate heat treatment can mitigate residual stresses, convert the initial martensitic microstructure into a stable α+β phase, and enhance ductility while maintaining strength. PERUMAL et al. [¹⁸] studied the impact of solution treatment with short- and long-term aging to evaluate the mechanical, wear and corrosion properties of Grade 5 Ti alloy. The alloy was subjected to solution treatment at 960 °C for 60 minutes, in two distinct aging conditions: short-term ageing (DHTS), for 120 s and long-term aging (DHTL), for 300 min at a temperature of 550 °C. The experimental findings reveal that short-term (DHTS) ageing improved mechanical, wear, and corrosion properties. Short-term aged samples showed a 25.61% increase in hardness and a 23.26% higher tensile strength compared to the as-cast alloy. The DHTS sample exhibited the lowest wear rate, 3.4 times lower than the as-cast and 1.6 times lower than the long-term aged counterpart (DHTL). The significant improvement in DHTS samples was due to the presence of α’– α’-martensite phase and β phase that evolved during quenching and heat treatment conditions. QIANGQIANG et al. [¹⁹] investigated the effect of solution and ageing heat treatments on the microstructure and mechanical performance of hot-rolled Ti–6Al–4V alloy sheets. They noticed the transformation from α → β → β_t. This led to the transformation from α-phase elongated structures to equiaxed structures, alongside a growth in β-phase content. The resulting bimodal microstructure, particularly rich in β-transformed (β_t) regions with dense nano-needle α structures, significantly enhanced strength. However, excessive β_t content introduced brittleness. Optimal strength-ductility balance was achieved at 930 °C for 60 minutes or 960 °C for 30 minutes, where β_p and β_t phases existed in nearly equal proportions, maximizing performance stability.

BALTATU et al. [²⁰] studied the effect of heat treatment for improving the properties of a few titanium- based alloys intended for biomedical use. The heat treatment cycle consisted of a high temperature quenching accomplished in three steps (650 °C for 25 min, 850 °C for 20 min and 950 °C for 20 min). The heat treatment included multi-stage high-temperature quenching followed by gas-assisted cooling and tempering, designed to improve microstructure and mechanical performance. Optical microscopy revealed the notable structural transformations that include acicular and coarse structures found in the case of β-Ti alloys. Micro hardness exhibited improvement in some alloys, such as Ti15Mo7Zr15Ta1Si, which showed approximately 5% increases, while others, like Ti15Mo7Zr15Ta0.5Si, saw reductions up to 9%. Moreover, indentation testing indicated that the modulus of elasticity ranged between 36.25 and 66.24 GPa, well within the desirable range for biomedical applications. Jin-Wen et al. [²¹] investigated the effects of changing Mo content on the microstructure and β-grain growth behaviour of Ti–Mo binary alloys. They discovered that refined grain structures resulted from higher Mo levels (>10 wt%), which fully stabilised the β phase and inner acicular α^’. Remarkably, the Ti–20Mo alloy depicts only retained β phase disp4layed slower growth and a significantly higher activation energy (272.16 kJ/mol) than the Ti–4Mo alloy, which displayed faster grain growth and a lower activation energy (83.3 kJ/mol). This was explained by the “solute drag effect,” which suppresses grain coarsening by preventing atomic mobility at grain boundaries due to Mo atoms. Grain refinement and solid solution strengthening were the main mechanical effects of higher Mo content, which improved room-temperature strength and plasticity. KOTOV, et al. [²²] investigated the increase on the molybdenum (Mo) content of 1 to 5 wt% improved the super plasticity of Ti–4Al–1V–1Fe–1Ni–0.1B alloy. They showed that Mo significantly impacts microstructural evolution and mechanical performance, particularly at lower deformation temperatures, by adding different Mo contents (1–5 wt%). Mo’s effect was most noticeable at 625 °C, even though high-temperature forming (700–775 °C) produced excellent superplastic behaviour across all compositions with elongation-to-failure up to 1000%. Increasing Mo improved both strain rate sensitivity (m = 0.5) and elongation (~700%) by raising the β-phase content above the critical 20%. Furthermore, room-temperature strength and ductility were increased by 5% Mo content, demonstrating the effectiveness of Mo as a slow-diffusing β-stabilizer that inhibits grain coarsening and enhances structural integrity after forming.

The aim of the present study is to explore the possibilities of recommending Ti-6Al-5Zr-0.5Mo-0.2Si alloy as a potential material in replacement to conventional Ti–6Al–4V alloy, which is widely used in aerospace applications. It is observed that the research work proposed on Ti-6Al-5Zr-0.5Mo-0.2Si alloy is seldom found in the literatures. Studies pertaining to dynamic strain aging and silicide precipitation as the two major investigations carried to evaluate the performance characteristics of the Ti-6Al-5Zr-0.5Mo-0.2Si alloy [²³]. For applications such as compressor blades, it is desired to ensure materials having stable elevated-temperature properties, reinforced by α^’-rich microstructures that provide higher performance, excellent creep resistance, thermal stability, oxidation resistance, and high resistance to fatigue crack growth. Ti-6Al-5Zr-0.5Mo-0.2Si, a less-studied near-α titanium alloy is not investigated in this regard for the extreme temperatures [²⁴]. As a result, the heat treatment of the Ti-6Al-5Zr-0.5Mo-0.2Si alloy to near the beta trans temperature can impart not only better strength but also improve the toughness, creep resistance and inhibit fatigue crack propagation in the compressor blade when subjected to severe harsh environmental conditions [²⁵]. This work provides a multi-dimensional overview on studying the effect of solutioning and ageing heat treatments of Ti-6Al-5Zr-0.5Mo-0.2Si alloy on the microstructural evolution, mechanical and high temperatures wear behaviour properties.

2. EXPERIMENTAL PROCEDURE

For the present study, Ti-6Al-5Zr-0.5Mo-0.2Si alloy was chosen as the base material. The base material was prepared from ingots (initially fabricated using a vacuum arc remelting process) and later forged at 1000 °C, below the β-trans temperature, to produce it as a billet. Later, the extrusion process transformed the billets into a round bar of 20 mm in diameter and 170 mm in length. The beta trans temperature of the as-received base material was measured using the metallographic method and found to be 990 ± 5 °C. The elemental composition (by energy dispersive spectroscopy) and mechanical properties of the base material Ti-6Al-5Zr-0.5Mo-0.2Si alloy were presented in Tables 1 and 2.

Table 1
Elemental composition of the base Ti-6Al-5Zr-0.5Mo-0.2Si titanium alloy.

ALLOY	COMPOSITION
ELEMENTS	Ti	Al	Zr	Mo	Si	Fe	O	C
Wt %	Bal	6.28	5.23	0.51	0.23	0.02	0.02	0.04

Table 2
Mechanical properties of Ti-6Al-5Zr-0.5Mo-0.2Si titanium alloy in the as received condition.

DENSITY g/cm3	YOUNG’S MODULUS GPa	ULTIMATE TENSILE STRENGTH (UTS), MPa	YIELD STRENGTH (YS), MPa	PERCENTAGE ELONGATION (PE), %	VICKERS HARDNESS, HV
4.5	110	984 ± 6	834 ± 5	14	315 ± 5

2.1. Process of heat treatment

To enhance the performance attributes of the Ti-6Al-5Zr-0.5Mo-0.2Si alloy, the heat treatment was performed in a vacuum furnace under a pressure of approximately 10^-5 MPa. The specimens underwent solution heat treatment at 850°C, 950°C, and 1050°C for 1 hour, followed by oil cooling after the solutionizing process. Later, the ageing heat treatment was performed at 600°C for a 4-hour duration for the specimens of 850°C, 950°C and 1050°C and was further air cooled to the room temperature [²⁶]. Figure 1 illustrates a schematic overview of the solution treatment and aging procedure implemented for the as-received titanium alloy.

Figure 1
Heat treatment cycle designed to enhance the performance of Ti-6Al-5Zr-0.5Mo-0.2Si.

2.2. Metallographic analysis

Microstructural analysis was carried out using scanning electron microscopy (SEM) to examine the morphological features and phase distribution within the alloy. The specimens were prepared following conventional metallographic procedures, including polishing and etching, to ensure proper surface preparation for microscopic examination. Polishing was followed by Kroll’s reagent, which was used to etch the samples and reveal the underlying microstructure. Phase identification and analysis of possible transformations were carried out using X-ray diffraction (XRD) with a Philips X’pert PRO system, utilising CuKα radiation at an operating voltage of 40 kV and a current of 40 mA during the X-ray diffraction analysis.

2.3. Mechanical behavior

Two specimens were extracted from each round bar and prepared according to the ASTM E8-04 sub-size standard to assess yield strength, ultimate tensile strength, and elongation, which were measured to evaluate the mechanical performance of the alloy. Tensile tests were carried out using a computerised universal testing machine at a crosshead speed of 2 mm/min, and the average values were documented. Microhardness measurements were executed using a Vickers micro hardness tester, applying a 0.5 kgf load with a dwell time of 10 seconds.

2.4. Tribological studies

Wear testing of the different heat-treated specimens was performed using a pin-on-disc apparatus (Make: Ducom; Model: TR-20LE) to evaluate their wear resistance under controlled conditions. The tests were conducted under two conditions: room temperature (30 °C) and elevated temperature (200 °C). The specimens were prepared according to the ASTM G99 standard, using a pin with a 6 mm diameter. Wear tests were carried out against commercial GCr15 steel discs, each 8 mm thick and 70 mm in diameter. The GCr15 steel was austenitized at 900°C, oil quenched, and then tempered at 400°C for 2 hours to achieve a hardness of 56 HRC. In the experimental setup, a track diameter of 100 mm and a sliding distance of 750 mm were maintained. Each test was conducted over a total duration of 2400 seconds. The applied load, sliding velocity, and disk rotational speed were kept constant at 20 N, 2.68 m/s, and 640 rpm, respectively. Prior to every test, the contact surface of the pin was polished using a 400-grit SiC emery paper to attain a surface roughness of around 0.4 µm, followed by cleaning with ethanol. The mass loss in the specimens were determined using an electronic balance with a precision of 0.01. Similarly, the wear loss and coefficient of friction of each sample were determined for three consecutive specimens and their standard deviation values are presented.

3. RESULTS AND DISCUSSIONS

3.1. Microstructural examination

Fig. 2(a) depicts the microstructure of the as-received Ti-6Al-5Zr-0.5Mo-0.2Si. The α grains are heavily deformed, resulting in hot forged and annealed heat treatment conditions. The primary α-grains are platelet-shaped. However, due to the swaging effect, some grains had transformed to equiaxed α grains as seen in Figure 2(a). The average grain size of the as-received base metal was measured using the Mean Linear Intercept method and was found to be 50 µm; however, the variation in the grain size was found to be between the range of 35–65µm.

Figure 2
SEM microstructure of as-received material (a) Ti-6Al-5Zr-0.5Mo-0.2Si alloy & solution treated and aged specimens (b) 850°C (c) 950°C (d) 1050°C.

Fig. 2b, 2c, 2d illustrate the microstructure of the samples subjected to solution-treated conditions of 850°C, 950°C and 1050°C for 1 hour, followed by ageing heat treatment at 560°C for a four-hour duration. It was inferred that the microstructural phase was significantly unchanged at 850°C, as the heat-treated temperatures are less than the β trans temperature of the as-received base metal. However, a significant increase in the primary α-α-grain size with dispersed β phase was observed, as depicted in Fig. 2(b). Fig. 2(c) illustrates the evolution of β-phase microstructure widespread across the primary α-grain boundaries. The fraction of β phase microstructure was slightly higher than that of the specimens of 850°C, as the heat treatment temperature is almost near the β trans temperature of the as-received base metal. In Ti alloys, the β-phase can form via solid-state phase transformations, such as eutectoid or eutectic reactions, or directional solidification. The β phase has lamellar structures and signifies that when the material is heated above the β- trans temperature, and then rapidly cooled by oil quenching, it leads to an increase in the fraction of retained β phase microstructure, which is metastable. Reheating at modest temperatures through aging speeds up the phase transition by adding thermal energy causing islands of beta phase transform to secondary alpha precipitates at the interface of α/β with high interfacial energy [²⁷]. This tends to harden and strengthen the ability of the material. At 1050°C, the primary α grain boundaries were transformed entirely to β- phase, however, the during the ageing process, the β- phase transformed to the secondary α -grain boundaries forming acicular structure, owing to lower concentration of β-phase contributing elements such as molybdenum [²⁸]. The microstructural arrangement comprises secondary alpha grain boundaries and beta colony as seen in Fig. 2 (d).

The results of XRD analysis also substantiates the above fact as depicted in Fig. 3.

Figure 3
XRD analysis.

Table 3 illustrates the concentration of molybdenum and silicon in the matrix of Ti-6Al-5Zr-0.5Mo-0.2Si alloy at different heat treatment temperatures. As the heat treatment temperature increased, the concentration of molybdenum in the matrix decreased, promoting a reduced fraction of the β phase within the matrix of secondary α grains. In contrast, the silicon content gradually increased with rising heat treatment temperatures.

Table 3
Elemental composition of molybdenum and silicon.

SAMPLE	Wt. % OF Mo	Wt. % OF Si
Base metal	0.51	0.23
850℃	0.45	0.25
950℃	0.39	0.29
1050℃	0.31	0.29

Silicide precipitation is a critical aspect in titanium alloys, occurring when silicon (Si) interacts with titanium or other alloying elements to form titanium silicide (Ti-Si) phases. Ti₅Si₃ is the most common precipitation that affects the mechanical properties. Ti₅Si₃ is a complex and brittle intermetallic phase, which forms at temperatures greater than 600 °C during ageing heat treatments. The precipitation of Ti₅Si₃ causes embrittlement at the grain boundaries and promotes the crack initiation [²⁹]. In the case of Ti-6Al-5Zr-0.5Mo-0.2Si alloy, silicide precipitation occurs more slowly than in other alloys, primarily due to its higher aluminium content and relatively low silicon concentration. Aging the alloy at 550°C for one week resulted in limited silicide formation, primarily along grain boundaries. At the same time, the standard commercial heat treatment did not produce any clearly detectable silicide particles [³⁰]. FLOWER et al. [³¹] noted that standard heat treatments applied to near-α titanium alloys containing up to 0.5% silicon typically do not lead to the formation of silicide dispersions. However, noticeable improvements in strength and creep resistance were observed, likely attributed to the segregation of silicon at grain boundaries, which may act to hinder dislocation movement. It was also confirmed from XRD that, aside from the primary α, β phases and secondary α precipitates, no secondary phases were observed in the heat-treated and aged samples in the temperature range of 850°C to 1050°C.

3.2. Evaluation of mechanical properties

Fig. 4 illustrates the tensile properties of the Ti-6Al-5Zr-0.5Mo-0.2Si alloy subjected to different heat treatment temperatures. The yield strength and the ultimate tensile strength of the as-received Ti-6Al-5Zr-0.5Mo-0.2Si alloy are 834 ± 5 MPa and 980 ± 6 MPa. The yield strength and ultimate tensile strength of heat-treated specimens at temperatures of 850 °C, 950 °C, and 1050 °C are 893 ± 5 MPa and 1076 ± 6 MPa, 900 ± 7 MPa and 1126 ± 9 MPa and 920 ± 8 MPa and 1270 ± 10 MPa. With an increase in temperature, the strength of the heat-treated Ti alloys increased significantly, with a decrease in ductility. Fig. 5 depicts the tensile test specimens prepared as per the ASTM standards. Fig. 5 (a) and (b) illustrate the samples before and after testing conditions. The elongation percentage of the as-received base metal was 14% and for heat-treated specimens at temperatures of 850 °C, 950 °C, and 1050 °C were 12%, 9% and 7% respectively.

Figure 4
Evaluation of tensile properties.

Figure 5
Tensile test specimens (a) before test (b) after test.

Furthermore, as the heat treatment temperature increased, the strength of the titanium alloy improved, complemented by a noticeable decrease in ductility. The improvement in strength was due to the presence of lower β phase stabilisation across the primary α-grain boundaries through solution strengthening with the presence of elements such as molybdenum and zirconium particles at the grain boundaries, which inhibited the dislocation movement in the case of 850 °C. The highest tensile strength was achieved in the case of 1050 °C, which was due to the transformation strengthening mechanism, where the transformation from β to secondary α precipitates resulted in a fine secondary α microstructure within the vicinity of the retained β phase. Thus, the finer secondary α-microstructure significantly improved the strength of the heat-treated Ti alloy. The microhardness results are also depicted in Fig. 6, which illustrates a similar trend to the strength factor. It is observed that the increase in the microhardness is due to the evolution of more secondary α grains which emerged during aging at elevated heat treatment temperatures.

Figure 6
Evaluation of microhardness.

The observed decrease in ductility of the heat-treated specimens can be associated with the slip transmission mechanism, which governs the dislocations movement across phase boundaries within the microstructure [³², ³³]. A slip transmission mechanism is considered adequate when the crystallographic orientation between the β and α phases satisfies the conditions defined by the Burgers orientation relationship, either fully or partially. Under these circumstances, the slip is successfully transmitted from the α phase into the neighbourhood of the β phase. However, if the Burger’s relationship is not met, the slip transferred from α-phase is restricted and fails to pass through the interface of α/β. Thus, the stress concentration produces deformation mismatch or incompatibility. This hinders slip transmission; hence, a remarkable strength is achieved with lower ductility. As a result, due to the increase in the heat-treated temperatures of the Ti alloy, there is a considerable decrease in ductility. The decrease in the ductility of the heat-treated specimens can be further validated with the SEM fractography. The fractured surfaces were observed to depict more fibrous features in the case of as-received-based metal and heat-treated specimens at 850 °C, seen in Fig. 7(a) and (b). While at 950 °C and 1050 °C, more facets are observed at the fracture location, due to a hindered slip transmission mechanism as seen in Fig. 7(c) and (d).

Figure 7
Tensile fracture characteristics of the specimens tested (a) base material, (b) 850 ℃, (c) 950 ℃, (d) 1050 ℃.

3.3. Wear analysis

Figure 8 (a–b) illustrates the coefficient of friction of the samples tested at room temperature (30 °C) and elevated temperature (200 °C). At the same time, Fig. 8(c) displays the overall wear rate of Ti-6Al-5Zr-0.5Mo-0.2Si alloy and the heat-treated specimens subjected to at room temperature (30^oC) and elevated temperature (200^oC) dry sliding testing conditions. At room temperature, the friction coefficient and wear rate for the as-received alloy were 0.3327 and 5.4 × 10⁻³ mm³/Nm, respectively. At the same time, the heat-treated samples exhibited an increased friction coefficient of 0.3502 and a wear rate of 5.76 × 10⁻³ mm³/Nm, 0.3721 and a wear rate of 5.32 × 10⁻³ mm³/Nm and 0.4202 and a wear rate of 4.83 × 10⁻³ mm³/Nm at 850 °C, 950 °C, and 1050 °C, respectively.

Figure 8
Coefficient of friction and wear rate of Ti-6Al-5Zr-0.5Mo-0.2Si alloy at different heat-treated specimens at varying temperatures (a) room temperature 30℃ (b) 200℃ (c) wear rate of different heat-treated samples.

Similarly, at 200 °C, the friction coefficient and wear rate for the as-received Ti-6Al-5Zr-0.5Mo-0.2Si were 0.4227 and 5.82 × 10⁻³ mm³/Nm, respectively. At the same time, the heat-treated samples exhibited an increased friction coefficient of 0.4202 and a wear rate of 5.23 × 10⁻³ mm³/Nm, 0.4352 and a wear rate of 5.13 × 10⁻³ mm³/Nm and 0.4402 and a wear rate of 4.76 × 10⁻³ mm³/Nm at 850 °C, 950 °C, and 1050 °C. It is observed that the base material is subjected to severe wear at 200 °C testing conditions when compared to the base samples tested at room temperature conditions. In contrast, the wear resistance of heat-treated samples when tested at 200 °C is comparatively better than that of the heat-treated samples tested at room temperature. Figure 9 presents the graphical representation of wear loss in Ti-6Al-5Zr-0.5Mo-0.2Si over time at different heat treatment temperatures, 30°C (room temperature) and 200°C. The wear behaviour at 30 °C and 200 °C depicts almost a similar trend of wear behaviour; however, the magnitude of wear at 200 °C was lower than at 30 °C. In both environments, wear depth increased progressively with sliding time; however, the heat-treated specimens showed significantly reduced wear compared to the base material. The 1050 °C-treated sample exhibited the lowest wear rate, attributed to enhanced microstructural stability and development of a possible defensive tribo-oxide layer [³⁴, ³⁵]. At 200 °C, the wear resistance of all specimens improved slightly, indicating a temperature-assisted oxidative wear mechanism becoming more dominant, especially in higher- temperature-treated specimens. Figs. 10 and 11 illustrate the SEM wear morphologies of the specimens tested at room temperature and high temperature (200 °C) conditions. It is observed that a minimum wear was obtained in the case of 1050 °C, followed by 950 °C, 850 °C, and the as-received sample when tested at room temperature conditions. However, a slightly lesser wear was observed in the specimens when tested at 200 °C, following a similar trend line as the room temperature conditions.

Figure 9
Wear behaviour of Ti-6Al-5Zr-0.5Mo-0.2Si alloy at different heat-treated specimens at varying temperatures (a) 30℃ (b) 200℃.

Figure 10
SEM wear morphologies of the specimens tested at room temperature (30°C) (a) ARS, (b) 850°C, (c) 950°C, (d) 1050°C.

Figure 11
SEM wear morphologies of the specimens tested at high temperature (200°C) (a) ARS, (b) 850°C, (c) 950°C, (d) 1050°C.

At room temperature conditions, more wear debris and deeper groove cuts are predominant in the case of specimens of as-received and 850°C. The lower hardness and unfavourable grain sizes effects due to lower heat treatment temperatures than the β-trans temperatures might be the reason for both abrasive and adhesive wear mechanisms [³⁶, ³⁷]. At room temperature, protective oxide films form more slowly and are less stable, resulting in higher friction coefficients, contributing to poor wear resistance. However, at 950°C and 1050°C, the wear resistance has been improved due to microstructural modification with β-transformed grains and the presence of secondary α precipitates that provide superior wear resistance and better work hardening capabilities, improving resistance to plastic deformation during wear. As a result, abrasive wear is more dominant at 950°C and 1050°C specimens.

At 200 °C, there is a considerable improvement in wear resistance in the heat-treated specimens when compared with the results of the room temperature test specimens. Under these temperatures, oxidation becomes more pronounced, especially in Ti alloys. This leads to the development of a stable and adherent oxide layer primarily composed of TiO₂ along with oxides of Al and Zr which serves as a protective barrier, minimizing direct metal-to-metal contact. This results in reducing the friction and material removal, as a result, the transitions in the mechanism of wear are governed by oxidative or tribo-oxidative wear, which is less severe than abrasive wear [³⁸]. While some level of abrasive wear still occurs (as oxides can break off and act as third-body abrasives), the overall wear rate is lower at 200°C because the beneficial effects of the stable tribo-oxide layer outweigh any negative abrasive processes.The results of EDS are shown in Figs. 12 & 13 also substantiate the above fact. A typical EDS of the heat-treated specimens at 1050^oC is illustrated in Fig. 12, which confirms the oxygen pickup in the wear surface of sample tested at 200^oC while less oxygen pickup compared to the same sample tested at room temperature conditions in Fig. 13. To conclude, the superior wear resistance 1050°C treated specimens’ likely results from their ability to form and maintain a stable triboxide layer, while the as- received and 850°C specimens may experience more adhesive wear and abrasive layer formation.

Figure 12
EDS analysis of the worn-out sample 1050℃ tested at 200℃.

Figure 13
EDS analysis of the worn-out sample 1050℃ tested at room temperature conditions.

4. CONCLUSIONS

1. The heat treatment process demonstrated a significant enhancement in the mechanical properties of the Ti-6Al-5Zr-0.5Mo-0.2Si alloy.

2. Solution heat treatment and ageing at a temperature slightly higher than β trans- temperature resulted in the formation of fine secondary α microstructure within the vicinity of the retained β phase.

3. The highest tensile strength of 1270 ± 10 MPa, 29% more than the base material, was achieved in the case of 1050°C due to the transformation strengthening mechanism, which exhibited excellent strength and superior wear resistance.

4. A higher wear rate resistance of 4.39 × 10⁻³ mm³/Nm was achieved in the heat-treated sample at 1050 °C when tested at 200 °C, due to the presence of a mild, durable and stable tribo oxide layer.

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Publication Dates

• Publication in this collection
  14 Nov 2025
• Date of issue
  2025

History

• Received
  04 June 2025
• Accepted
  24 Sept 2025