Revealing High-Temperature Oxidation and Tensile Behaviors along with Underlying Mechanisms of a Titanium Alloy with Precipitated Titanium Silicide

Abstract

Titanium alloys, with their impressive strength relative to their weight, resistance to corrosion, and compatibility with biological systems, have found extensive applications in various industries. In high-temperature environments, especially within the aerospace sector, it is essential to advance titanium alloys that boast enhanced resistance to oxidation and superior mechanical characteristics. This work investigates the oxidation characteristics and mechanical performances at high temperatures of a titanium alloy with titanium silicide particles. Oxidation at temperatures of 600–700 °C over a span of 8–32 h led to the formation of protective oxide layers and moderate oxidation rates. However, accelerated oxidation and oxide spallation occurred after exposed at 800 °C for a period of 16 h, indicating inadequate oxidation resistance over 800 °C. Subsequent tensile tests at 650 °C revealed intricate dislocation patterns in the α-Ti matrix and their strong interaction with interfaces of α-Ti/Ti₅Si₃, which is indicative of an efficient load transfer between the precipitates and the matrix. Overall, this study offers fresh perspectives on the oxidation kinetics and the deformation processes of titanium alloys with in-situ Ti₅Si₃ particles at high temperatures. These insights will guide subsequent alloy development endeavors aiming to broaden the use of titanium alloys in increasingly challenging high-temperature settings.

1. Introduction

Due to their standout features, like their lightweight nature, superior strength, impressive corrosion resistance, biocompatibility, and absence of magnetic properties, titanium alloys have become indispensable in diverse industries [1]. These attributes have rendered titanium alloys indispensable in critical sectors, such as in aerospace, military, marine, automotive, and industrial applications [2]. In particular, the unique demands of titanium alloy applications under high-temperature conditions have prompted the need for innovations in the design of materials. Notably, the aerospace domain necessitates materials capable of enduring elevated temperatures and sustained stresses, making titanium alloys with an enhanced high-temperature performance vital for next-generation aero-engines and aircraft [3,4]. Consequently, efforts to enhance their high-temperature oxidation resistance and mechanical properties have become a focus of extensive research.

However, a challenge arises when titanium alloys encounter oxygen environments exceeding 600 °C; there is a risk of forming unstable oxides, with a subsequent detachment from the surface [5]. To address this issue, extensive research has been dedicated to introducing ceramic particles with improved high-temperature durability, high modulus, and oxidation resistance [6,7,8]. It was found that TiB whiskers diminished the oxidation resistance in the titanium alloy, while Ti₅Si₃ particles bolstered the oxidation resistance by refining TiO₂ particles, thereby enhancing the toughness and compactness of the oxide layer [9]. The addition of LaB₆ significantly decreased the oxidation rate of titanium alloys at 600 °C [10]. The parabolic rate constants of the Ti-Al-Sn-Zr alloy at 750 °C and 850 °C decreased with the increase in the added SiC particles [11]. On the other hand, various titanium alloys incorporated with ceramic particles, such as TiC, TiB, NbC, Ti₃AlC, and La₂O₃, have been shown to bolster their mechanical characteristics, particularly by elevating their flow stress [12,13,14,15,16]. Moreover, high strength–ductility synergy is highly required for titanium alloys with ceramic particles, both at room temperature and at elevated temperatures. This balance is vital to cut down production expenses and mitigate the potential for failures during usage.

During high-temperature deformation processes in particular, the grain boundary slip mechanism takes precedence. The grain boundaries are susceptible to microcrack formation due to the significantly reduced strength at elevated temperatures, and its rupture could lead directly to the onset and spread of microcracks [17,18]. Consequently, in comparison with the alloys reinforced with foreign particles from outside, titanium alloys with strengthening ceramic particles through in-situ precipitation have the advantages of fine reinforcements, clean reinforcement/matrix interfaces with no contamination, superior wettability and interfacial bonding, as well as low preparation cost [19]. In addition, silicide ceramics demonstrated better deformation characteristics than other compounds like carbides and borides [20]. Specifically, the silicide of Ti₅Si₃ exhibits a high melting point, low density, pronounced creep resistance, high fracture toughness over 1000 °C, good oxidation resistant behavior, and a thermal expansion coefficient closely matching that of titanium [21,22,23]. In this context, this work presents the investigation of a titanium alloy incorporating titanium silicide precipitates. Given the material’s potential, this study focuses on understanding its oxidation behavior and mechanical response at high temperatures. By systematically exploring the interplay of in-situ Ti₅Si₃ precipitates and the alloy matrix, this investigation contributes to advancing the comprehension of high-temperature alloy behavior and paves the way for novel applications in aerospace and other demanding environments.

2. Materials and Methods

In this research, we investigated a high-strength titanium alloy with the following composition: Ti-4.5Al-2.5Si-2Mo-2Zr-2Nb-1Sn-1Ta-1W-1Hf (wt.%). The initial step involved melting the respective raw elemental materials 3–4 times in a vacuum arc furnace. Subsequently, the billet underwent hot forging at 1050 °C, followed by hot rolling at 850 °C. To analyze the phase composition, the X-ray diffraction (XRD) analysis was conducted on a 1 cm × 1 cm × 1 cm sample using a 7000S X-ray diffractometer (XRD, Shimadzu, Kyoto, Japan), employing Cu Kα radiation (wavelength λ = 1.5418 Å) at an operational setting of 30 kV and 30 mA. The scans covered a 2θ range, extending from 20° to 80°, utilizing a step size of 0.02°. The scanning proceeded at a speed of 1° per minute. To enhance the diffraction intensity, the sample was set to rotate at a rate of 30 revolutions per minute. An FEI G20 Transmission Electron Microscope (TEM, Thermo Fisher Scientific, Waltham, MA, USA) further validated the microstructure. For transmission electron microscopy (TEM) sample preparation, electrochemical thinning was employed. This method, also known as electropolishing, provides several advantages over traditional ion thinning techniques, particularly for materials like titanium alloys that can be challenging to prepare due to their unique chemical and physical properties. Electrochemical thinning offers a gentler approach that minimizes heat and stress effects on the sample, allows precise control over the thickness, and often enables faster preparation. This was meticulously carried out using an 8% perchloric acid solution in ethanol. The procedure was executed at a controlled voltage of 50 V, with the temperature meticulously maintained at 35 °C, ensuring optimal conditions for sample preparation. In addition, titanium alloy plates were fashioned into 8 mm × 8 mm × 8 mm squares using wire electric discharge. These specimens were then ground, cleaned, and prepped for oxidation tests. Oxidation tests were conducted in a tubular furnace, with static air as the controlled environment. Specimens were put into the furnace and exposed to temperatures of 600 °C, 700 °C, and 800 °C. Following specific durations (8 h, 16 h, 24 h, and 32 h), they were allowed to cool by air until room temperature. An electronic balance, accurate to 0.0001 g, recorded the sample mass change. The below formula was employed to compute the oxidative mass gain for each sample pre- and post-oxidation:

$$\Delta w=\frac{\Delta m}{S}$$

(1)

Here, Δw delineates the weight increase due to oxidative per unit surface area, quantified as mg/cm². The term Δm signifies the differential mass of the specimens pre- and post-oxidation, expressed in mg; S demarcates the aggregate surface area of the given specimen, measured in cm². Subsequently, the oxidation progression was visualized by mapping the kinetic curves between Δw and the time of t. In alignment with Wagner’s oxidation paradigm, the weight augmentation Δw is concomitant with the duration of oxidation t, as expounded in reference [24].

$${\left(\Delta w\right)}^n=k_{p}t$$

(2)

where Δw signifies the increase in mass from oxidation per unit of surface area; n denotes the index of the reaction; k_p stands for the constant of the reaction rate; and t designates the time period of oxidation. Additionally, the reliance of the temperature on k_p can be calculated according to the well-known Arrhenius formula [25].

The samples’ overall appearances and element distribution after undergoing oxidation were examined using a JEOL-6700 scanning electron microscope (SEM, JEOL, Tokyo, Japan), supplemented with the energy dispersive X-ray spectroscopy (EDX). Before SEM analysis, the samples were grinded and polished with the following etch in a chemical mixture (HF: HNO₃: H₂O at a ratio of 1:2:5 by volume) for a duration of 10–15 s. Subsequently, they were ultrasonically cleaned in both alcohol and water. After preheating for 5 min at a specified temperature, the tensile properties were evaluated at a steady strain rate of 5 × 10⁻⁴ s⁻¹ using an Instron-3382 at elevated temperatures of 550 °C, 600 °C, and 650 °C.

3. Results and Discussions

Figure 1a displays the XRD profiles of the titanium alloy, characterized by the presence of the α phase, β phase, and a minor quantity of the Ti₅Si₃ phase. A more in-depth view into the microstructural details can be seen in Figure 1b–o, which includes bright-field, STEM elemental distribution illustrations using EDX and SAED data. Figure 1m–o, obtained from the three blue regions indicated in Figure 1b, has been precisely indexed as α-Ti, β-Ti, and Ti₅Si₃, respectively, providing clear evidence of the crystalline structures present in the sample. More specifically, hexagonal structured Ti₅Si₃ was indexed according to both the interplanar spacing and angle verification to be the 2D projection from the zone axis of [10-1]. The interplanar spacings of the 010, 111, and 101 planes of this phase, as measured by the inverse of the lengths of the three sides from the respective spot to the transmission spot, match well with the standard values, with a tolerance of only 0.07 Å. Meanwhile, the tolerance in angles is found to be 1° only upon angle verification. The comparative reduction in Ti and the increase in Si made the morphology of the Ti5Si3 particles evident, as discerned in Figure 1d,f. Additionally, Figure 1c shows partial Hf dissolution into Ti₅Si₃, corroborated by an EDX quantitative analysis result of Ti_66.47Si_24.51Al_0.51Zr_0.57Mo_0.73Hf_4.33Ta_2.05W_0.83 in weight percentage. The unobvious concentration of Ta, W, and Mo (Figure 1k,l,g) in the spherical Ti₅Si₃ was attributed to their higher segregation in β-Ti; however, it is certain that the incorporation of Hf, Ta, W, and Mo would affect the properties of the silicide by altering its stability, diffusion kinetics, electron structure, and other performance factors. Previous research has indicated that the introduction of Hf enhanced the nucleation of fine Ti₅Si₃ particles, reducing the overall brittleness by leveraging its lower solution enthalpy in Ti₅Si₃ compared to both α-Ti and β-Ti [26,27]. STEM elemental distribution also revealed that particles A and B have much lower Si content compared to particle C, indicating that they are Ti-rich phases rather than silicides. As illustrated in Figure 1i,k,l, the images revealed an area enriched with β-Ti stabilizing elements, including Nb, Ta, and W, representing the β-Ti phase. In contrast, the surrounding area could be identified as the α-Ti phase.

For potential high temperature applications, understanding the oxidation behavior of the titanium alloy is crucial. The alloy samples underwent oxidation in open air at temperatures of 600 °C, 700 °C, or 800 °C for 8–32 h. The resulting mass change per unit area after oxidation is displayed in Figure 2a. It is evident that the rate of oxidation exhibited a gradual increase from 600 °C to 800 °C, with oxidation being slow at 600 °C and 700 °C, and rapid at 800 °C. In accordance with metal oxidation kinetics, the initial small weight increase at 800 °C for 8 h suggests the formation of a protective oxide film. However, as the duration of exposure increased, non-protective oxidation films accelerated the oxidation rate. By logging both sides of Equation (2) and graphing the variation of ln(Δm) against lnt, a linear trend emerged, as shown in Figure 2b. This allows for the extraction of the reaction index n and the rate constant k_p using linear regression. The obtained values were approximately 1.97 at 600 °C, 2.32 at 700 °C, and 1.69 at 800 °C for n, as well as 0.00714, 0.0775, and 0.4813 for k_p, respectively. The oxidation mechanism is predominantly controlled by diffusion processes. When n = 1, the oxidation reaction rate follows a linear law, governed by the oxidation reaction rate of oxygen and matrix metal [28]. For n = 2, the oxidation reaction rate follows a parabolic law, driven by reactant diffusion in the oxide film [29]. When n takes a value between 1 and 2, the oxidation process involves both reactant diffusion in the oxide film and oxidation reaction between oxygen and the base metal [30]. The parabolic rate constant k_p signifies the alloy oxidation rate. Figure 2 illustrates that n is 1.97 at 600 °C with a small k_p value, indicating excellent oxidation resistance. At 700 °C, the n value of 2.32 indicated a diffusion-controlled oxidation mechanism with effective substrate protection via the oxide film [31]. At 800 °C, an n value of smaller than 2 implies interface reaction control between oxygen and the oxide film/substrate, and a relatively large k_p value suggests that the oxidation process was potentially influenced by stress, pores, and grain boundaries [32]. As the oxidation temperature increased, the values of n decreased with the increasing k_p, reflecting combined diffusion and surface oxidation reaction control. Linear analysis, shown in Figure 2c, provided an oxidation activation energy value of 164.1 kJ/mol, which shed light on the oxide growth process and the diffusion of oxygen into the matrix [33].

Figure 3 depicts the overall color presentation of the alloy after oxidation at temperatures ranging between 600 and 800 °C for various time durations. It can be seen that, at 600 °C, the surface color of the alloy shifted from light-yellow to light-brown, eventually progressing to dark-yellow and culminating in a blue hue as oxidation proceeded. Similarly, oxidation at 700 °C resulted in color transitions from yellowish-brown to blue, while oxidation at 800 °C induced a change from blue to a gradually light-green shade. The intricate surface color variations of oxidized titanium alloys arise from factors including the oxide layer composition, layer compactness, thickness, interference effects of light, surface roughness, luminous flux of the oxide film, refractive index, and reflectance [34,35,36]. As the titanium alloys underwent surface oxidation, a diverse range of titanium oxides formed, each contributing distinct colors. Therefore, the uneven coloration of the oxide film can be attributed to the formation and content of these oxides. The porous oxide layer’s capacity to absorb light affected the interplay between light bouncing off the oxide film’s surface and the light reflecting at the oxide/titanium boundary. This leads to intricate visible hues characterized by varying wavelengths and proportions. It is worth noting that TiO₂ with crystalline irregularities can display a spectrum of colors like yellow, blue, and white [37,38].

Figure 4 displays the XRD patterns from the oxide layers that developed on the alloy surface after exposure to various temperatures and durations. When subjected to oxidation at 600 °C, a notable intensity in the Ti diffraction peak was observed. However, as the oxidation duration extended, there was a corresponding increase in the intensity of the TiO₂ peaks, accompanied by a gradual reduction in the Ti peak intensity. Conversely, at 800 °C, a significant decline in the Ti peak intensity was observed, coinciding with an elevation in the TiO₂ intensity. This trend signifies a marked decrease in the alloy’s resistance to oxidation. The observed variations in the oxide layer composition and content under distinct oxidation conditions align with the discernible differences in the oxide layer colors, as depicted in Figure 3. Importantly, the prevalence of TiO₂ on the alloy surface becomes evident in cases of oxidation at 800 °C.

Figure 5 depicts the surface view of the alloy after oxidation at temperatures between 600 °C and 800 °C for durations of 8 h, 16 h, 24 h, and 36 h, respectively. The oxide film formed on the surface was relatively dense and exhibited bright white characteristics (Figure 5a–d) at 600 °C, attributable to the lower oxidation temperature and partial preferential oxidation behavior occurring along the phase boundaries with high energy and diffusion rates. The oxidation reaction rate primarily governed the process at 600 °C. With the increase to the higher temperature of 700 °C (Figure 5e–h), a generally flat surface was observed, while the emergence of small holes was notable with the increasing oxidation time. This phenomenon might be attributed to the incomplete oxidation of positions occupied by oxidation-resistant Ti₅Si₃ particles during the process. These particles acted as barriers against further oxygen ingress, leading to the formation of small surface pores. The oxidation kinetics curve, which depicts a deceleration in the oxidation rate as the oxidation time increases, supports the above observation. The thicker oxide layer arising from the greater diffusion distance of the titanium ions curtailed the surface oxidation reaction rate, enabling a protective role. In contrast, oxidation at 800 °C yielded an uneven surface with nodules (Figure 5i–l). After 8 h of oxidation, a few pores persisted due to the high oxidation resistance of Ti₅Si₃. However, extended holding times resulted in the peeling off of oxide layers from the surface due to residual stress differences between the oxide layer and substrate arising from their different thermal expansion coefficients. Following 32 h of oxidation, oxidized particles coarsened, engendering a chaotic and loose surface structure with short rods and irregular pores, facilitating oxygen diffusion. The heightened oxide detachment indicated compromised high-temperature oxidation resistance, in accordance with previous oxidation kinetics analysis. EDX analysis on the alloy oxidized for 32 h at 800 °C corroborated the complete formation of TiO₂ on the surface. Therefore, Ti₅Si₃ demonstrates a dual role in the oxidation behavior and mechanisms of the alloy. It acted as a protective barrier at the lower temperatures of 600 °C and 700 °C, retarding oxidation through selective oxide formation. However, at the higher temperature of 800 °C, the dissolution of Ti₅Si₃ accelerated the oxidation process by altering the alloy’s microstructure and providing pathways for oxygen diffusion. In essence, the presence of Ti₅Si₃ particles in the alloy played a crucial role in its oxidation tendencies, shaping the emergence of protective oxide barriers and dictating the pace at which oxidations transpired. The nature of this influence varies at different temperatures, leading to distinctive mechanisms that are essential for comprehending the alloy’s oxidation behavior at elevated temperatures. Comparatively, as Ti-6Al-4V contains a toxic element of vanadium, the activation energy for oxygen oxidation was only 126 kJ/mol [39]. In addition, a typical high-temperature titanium alloy of Ti-6Al-7Nb consists of Al₂O, TiO, and TiO layers, in which the addition of 7 wt.%Nb and 1 wt.%W hindered the diffusion of oxygen, and the effect of Al on the oxygen absorption fostered the formation of a uniform and compact oxide film on the exterior [40]. Therefore, the present alloy showed relatively high oxidation resistance.

Figure 5 presents a surface view of the titanium alloy subjected to oxidation at 600–800 °C over 8 h, 16 h, 24 h, and 32 h (a–l), with spot analysis on the oxidated sample at 800 °C for 32 h (m).

The specific requirements of certain titanium alloys for applications in high-temperature environments, coupled with the inevitability of oxidation reactions during processing without a protective atmosphere, necessitate an examination of their high-temperature deformation behavior [41]. Here, we conducted tests and analyses to investigate the response of the alloy under atmospheric exposure. Figure 6a,b presents the engineering stress-strain data obtained from high-temperature tensile tests conducted at 550 °C, 600 °C, and 650 °C. At 550 °C, the tensile properties were characterized by a yield strength (σ_y) of 227.54 MPa, a tensile strength (σ_t) of 305.4 MPa, and a value of elongation (ε) of up to 107.7%. For the tests at 600 °C, the values were 113.52 MPa for σ_y, 142.58 MPa for σ_t, and 253.1% for ε. Similarly, at 650 °C, the tensile properties were 42.41 MPa for σ_y, 54.11 MPa for σ_t, and 694.7% for ε.

Figure 6c shows images of the alloy specimens used for the high-temperature tensile tests before and after failure at temperatures of 550 °C, 600 °C, and 650 °C. It is evident that the strength decreased as the deformation temperature increased, while its plasticity gradually increased. This occurrence can be considered to be a result of the diminished strength in both the grain and phase boundaries and the weakening of the solid solution strengthening effect [15,17]. During thermal deformation, the magnitude of flow stress hinges on the balance between dynamic softening and work hardening processes. Enhanced thermal activation at higher temperatures increases the frequency of atomic jump behavior and reduces the energy required for dislocations to surmount barriers [42]. As a result, the critical shear stress decreased, leading to lower flow stress. Meanwhile, thermally activated dynamic recovery and recrystallization processes occurred, consuming dislocations and decreasing flow stress [43]. The fracture morphologies of the alloy specimens after high-temperature tensile deformation are presented in Figure 6d–f. The occurrence of slip between the adjacent β and α colonies was heavily influenced by their crystal orientation, rendering them susceptible to crack nucleation. In general, stress concentration, caused by uncoordinated deformation at phase boundaries between the α, β, and Ti₅Si₃ phases, triggered crack nucleation. The specimens showcase a myriad of evenly distributed and delicate dimples, enriched with an abundance of tear edges, indicative of substantial plastic deformation capacity. It is noteworthy that the stress concentration during high-temperature stretching is not uniformly distributed, owing to the different thermal expansion coefficients for the three constituting phases. Specifically, at 600 °C and 650 °C with high deformations (253.1% and 694.7%), internal cracks emerged in some large Ti₅Si₃ particles (~10 μm), while no interface debonding occurred. This observation underscores the robust interface bonding strength between in-situ Ti₅Si₃ and the surrounding matrix.

The essence of high-temperature deformation for alloys is the interplay between dynamic hardening and softening mechanisms [17,44]. Dynamic hardening involves dislocation accumulation, the formation of locked dislocation configurations, and energy consumption for dislocation initiation. The dynamic softening mechanisms include dynamic recovery, dynamic recrystallization, and insulation heating [45]. Dynamic recovery facilitates the migration and reaction of dislocations, leading to reduced dislocation densities. Dynamic recrystallization involves the movement of recrystallized grain boundaries [46]. The plastic deformation of polycrystalline titanium alloys is predominantly governed by intragranular and grain boundary slips. The dynamics of α slip and the interplay between α/β phase slip existed within intragranular slip. Conversely, grain boundary movements incorporate both β grain boundary slip and α boundary slip [17]. The TEM analysis here sheds light on the resulting high strength–ductility synergy observed during high-temperature tensile deformation.

A complex dislocation structure was notably observed within the α-Ti phase post-deformation (Figure 7a,b). As the temperature increased, the critical resolved shear stress decreased sharply, leading to both pyramidal and basal slips during high-temperature deformation [47,48]. The reinforcing Ti₅Si₃ particles impeded excessive dislocation motion, resulting in dislocation tangles or pile-ups at the phase boundaries (Figure 7a,b). Elevated dynamic recovery occurred within the metastable β phase, and with the increased deformation temperature, both the intragranular and grain boundary strengths declined, albeit with a sharper reduction in the grain boundary strength. At 650 °C, the weakest region in the polycrystalline alloys was the grain boundary. Consequently, the mechanism of grain boundary slip took precedence during elevated temperature deformation, with the grain boundaries being prone to microcrack formation due to their inherent weakness [17]. For the phase boundaries between Ti₅Si₃ and the matrix, however, it can be observed that the in-situ precipitates displayed strong wettability and interfacial bonding.

A complex dislocation structure was notably observed within the α-Ti phase post-deformation (Figure 7a,b). The complex dislocation assemblies within the α-Ti phase (Figure 7a) indicated the occurrence of dynamic strain aging during hot deformation [47], which restricted dislocation motion through the diffusion of solute atoms to dislocation cores, enabling dislocation tangles and junctions [48]. The formed dislocation cell structures signified the dynamic balance between work hardening and softening [49]. As the temperature increased, the critical resolved shear stress decreased sharply, leading to extensive basal, pyramidal, and cross slips during high-temperature deformation [50,51]. The high stacking fault energy of α-Ti contributed to extensive cross slips near the peak strains [52]. The cross-slip intersection provides more deformation modes and facilitates dynamic recrystallization. Moreover, the dislocation pile-ups along the α/Ti₅Si₃ phase boundaries (Figure 7b) produce a substantially increased stress concentration that could promote crack initiation. Nevertheless, the excellent interface strength withstood the stress concentration without incurring debonding. This is due to the sound metallurgical bonding and lattice coherency in the in-situ synthesized phases. The reinforcing Ti₅Si₃ particles impeded excessive dislocation motion, resulting in dislocation tangles or pile-ups at phase boundaries (Figure 7a,b). Elevated dynamic recovery occurred within the metastable β phase, and with the increased deformation temperature, both the intragranular and grain boundary strengths declined, albeit with a sharper reduction in the grain boundary strength. At 650 °C, the weakest region in the polycrystalline alloys was the grain boundary. Consequently, the mechanism of grain boundary slip took precedence during elevated temperature deformation, with grain boundaries prone to microcrack formation due to their inherent weakness [17]. For the phase boundaries between Ti₅Si₃ and the matrix, however, it can be observed that in-situ precipitates displayed strong wettability and interfacial bonding. To summarize, the designed Ti alloy architecture optimized the synergistic properties at elevated temperatures via well-tailored phase constituents, interface characteristics, and dislocation mechanisms.

4. Conclusions

In conclusion, this study presents a titanium alloy containing in-situ titanium silicide particles, for which the elevated temperature oxidation and tensile behaviors were thoroughly investigated and characterized. The main findings can be summarized as follows:

• The alloy showed excellent resistance to oxidation at 600 °C and 700 °C. An oxidation reaction index (n) close to 2 and a small parabolic rate constant indicated favorable oxidation resistance. At 700 °C, the n value of 2.32 suggested that a diffusion mechanism governed the oxidation process, ensuring effective protection by the oxide film.
• At 700 °C, the presence of oxidation-resistant Ti₅Si₃ particles at the phase boundaries hindered their full reaction with oxygen, acting as barriers to oxygen ingress. Conversely, at 800 °C, the oxidation was impacted by a strong interfacial reaction between oxygen and the oxide film/substrate. The resulting oxide film exhibited a porous structure with TiO₂ rods and irregular pores, indicating a dramatic decline in oxidation resistance at an elevated temperature.
• The high temperature tensile data tested at 550 °C were σ_y = 227.54 MPa, σ_t = 305.4 MPa, ε = 107.7%; σ_y = 113.52 MPa, σ_t = 142.58 MPa, ε = 253.1% at 600 °C; σ_y = 42.41 MPa, σ_t = 54.11 MPa, ε = 694.7% at 650 °C. Notably, high deformations at 600 °C and 650 °C led to internal cracks in certain large Ti₅Si₃ particles; however, no interface debonding occurred due to the superior interface bonding strength with the matrix.
• Microstructural analysis post-deformation indicated the presence of a complex dislocation structure, primarily within the α-Ti phase. The increased deformation temperatures led to decreased intragranular and grain boundary strengths, with the grain boundaries becoming the weakest points in the alloy at 650 °C. The grain boundaries were prone to microcrack formation during high-temperature deformation due to their inherent weakness. In contrast, the interfaces between the Ti₅Si₃ precipitates and the matrix exhibited robust interfacial bonding strength.