High-strength and high-thermal-stability near-alpha titanium alloy and preparation method thereof

Abstract

The application discloses a kind of high-strength high-thermal stability near α Type titanium alloy and preparation method, belong to new alloy material design development field.The titanium alloy is by following mass percentage component composition:Al 5.5%~7.0%, Sn 5.5%~8.0%, Zr 5.0%~8.0%, Mo 0.5%~2.5%, V 0.5%~2.5%, remainder is Ti and inevitable impurity, and Zr / Sn=0.8~1.5.The application is based on electron concentration theory, by greatly improving Sn and Zr content, using Sn to improve electron concentration and promote α2 ordered phase precipitation to improve high temperature strength, using Zr has infinite solid solubility in Ti, and can form substitutional solid solution with Ti, produce solute drag effect and inhibit α2 phase coarsening to improve thermal stability, realize alloy high temperature strength and thermal stability synchronous promotion.Preparation method includes ingot smelting, breakdown forging and bar preparation.The alloy of the application reaches 1253MPa and 800MPa above respectively at room temperature and 600 DEG C tensile strength, is improved more than 19% and 30% respectively compared with Ti60 alloy, after 600 DEG C / 100h thermal exposure, room temperature elongation after fracture keeps 14.5% above, is suitable for aero-engine blade, disc piece and other components.

Description

Technical Field

[0001] This invention belongs to the field of new alloy material design and development, and in particular provides a high-strength, high-thermal-stability near-α-type titanium alloy and its preparation method. Background Technology

[0002] Near-alpha type high-temperature titanium alloys possess excellent specific strength, fracture toughness, and high-temperature mechanical properties, making them crucial materials for manufacturing key components of advanced aero-engines. Ti-1100, Ti60, and Ti600 alloys are representative high-performance near-alpha type high-temperature titanium alloys, capable of long-term operation at temperatures up to 600 °C. These alloys, manufactured under industrial conditions, exhibit a tensile strength of 950 Nm at room temperature. 1050 MPa, elongation after fracture 10% 15%; tensile strength at 600 ℃ is 550. 650 MPa, elongation after fracture 10% 20%. After heat exposure at 600 ℃ for 100 h, their tensile strength at room temperature is 950. 1100 MPa, while the elongation after fracture decreased to 5%. The thermal stability of near-α titanium alloys is a significant issue, accounting for 9%. New-generation aero-engines place higher demands on the thermal strength and stability of key components such as blades, disks, and impellers at 600 °C. Simultaneously improving the high-temperature strength and thermal stability of near-α titanium alloys at 600 °C is a crucial challenge for achieving high performance, lightweight design, and long service life in aero-engines.

[0003] Solid solution strengthening is the primary strengthening mechanism for near-α type titanium alloys. Al has a high solid solubility in α-Ti, which can significantly improve the alloy's room temperature and high temperature strength. Simultaneous addition of Al and Sn can form a primary solid solution, further enhancing the alloy's high temperature strength. When the content of Al and Sn exceeds the critical solid solubility in α-Ti, Ti3X (X = Al, Sn) compounds, i.e., the α2 phase, precipitate. During alloy deformation, the α2 phase can pin dislocations and strengthen grain boundaries, thus significantly improving the alloy's high temperature strength. However, during long-term high-temperature service, high volume fractions of the α2 phase are prone to coarsening and segregation, significantly reducing the alloy's thermal stability and plasticity. Existing research results indicate that the volume fraction, size, and distribution characteristics of the α2 phase can be controlled through compositional design, effectively refining the grains while achieving a uniform and dispersed distribution of the α2 phase. Adding rare earth elements such as Y and Sc to near-α type titanium alloys can effectively suppress the coarsening and segregation of the α2 phase during long-term heat exposure, thereby improving the alloy's thermal stability. However, rare earth elements have extremely low solid solubility in α-Ti and tend to segregate at grain boundaries or in localized areas, which has an adverse effect on the room temperature plasticity of the alloy.

[0004] Compositional design based on electron concentration theory is an important method for controlling the α2 phase in near-α type titanium alloys. In Ti-X (Al, Sn) alloys, as the electron concentration increases, the driving force for the transformation of the disordered α phase to the ordered α2 phase increases, promoting the nucleation and precipitation of the α2 phase, thereby improving the room temperature and high temperature strength of the alloy. However, when the electron concentration exceeds a certain critical value, the volume fraction of the α2 phase increases significantly, and it is prone to coarsening and segregation under long-term high-temperature thermal exposure conditions, leading to a weakening of dislocation pinning and stress concentration in local areas, reducing the strength, plasticity, and thermal stability of the alloy. How to design alloy systems with high electron concentration and effectively control the volume fraction, size, and distribution characteristics of the α2 phase is a significant challenge for simultaneously improving the high-temperature strength and thermal stability of the alloy.

[0005] Currently, high-temperature titanium alloys used at 600℃ and above generally employ a technical route of adding elements such as Nb, Ta, and W to the Ti-Al-Sn-Zr-Mo-Si basic system. This results in complex alloy systems (typically containing more than eight alloying elements) that rely on Si to form silicide strengthening phases to ensure creep resistance. However, while the addition of Si improves high-temperature creep resistance, it significantly reduces the alloy's thermal stability. Furthermore, the high density of Nb, Ta, and W elements increases the alloy's specific gravity and cost. How to simplify the alloy system and improve thermal stability while ensuring high-temperature strength is another pressing technical challenge in the field of near-α type titanium alloy research. Summary of the Invention

[0006] To address the above issues, this invention aims to simultaneously and significantly increase the content of Sn and Zr elements, utilize the synergistic effect of Sn and Zr elements, and precisely control the volume fraction, size, and distribution characteristics of the α2 phase, thereby simultaneously improving the thermal strength and thermal stability of near-α type titanium alloys. Through optimization of alloy element content and reasonable forging process, a high-strength and high-thermal-stability near-α type titanium alloy and its preparation method have been developed.

[0007] On one hand, the present invention provides a high-strength, high-thermal-stability near-α-type titanium alloy, wherein the mass percentage of each component of the near-α-type titanium alloy is 5.5 wt.%. 7.0 wt.% Al, 5.5 wt.% 8.0 wt.% Sn, 5.0 wt.% 8.0 wt.% Zr, 0.5 wt.% 2.5 wt.% Mo, 0.5 wt.% 2.5 wt.% V, with the balance being Ti and unavoidable impurities. The relationship between Zr and Sn was precisely controlled to be 0.8 ≤ Zr / Sn ≤ 1.5.

[0008] Al is the main α-stabilizing element in this alloy system, and its room temperature and high temperature strength are improved through solid solution strengthening. Considering both strength and ductility, its content is controlled at 5.5 wt.%. 7.0 wt.%, to avoid alloy embrittlement caused by excessive Al content.

[0009] Sn and Zr are the core elements that enable the performance breakthrough in this invention. Sn has a significant solid solution strengthening effect and can effectively promote the dispersed precipitation of the α2 phase by increasing the electron concentration, thereby improving high-temperature strength. Zr has infinite solid solubility in Ti and can form a substitute solid solution with Ti, which can inhibit the coarsening and segregation of the α2 phase at high temperatures and improve the thermal stability of the microstructure. This invention simultaneously increases the content of both to above 5.0 wt.%, and strictly limits 0.8 ≤ Zr / Sn ≤ 1.5, so that the precipitation-promoting effect of Sn and the stabilizing effect of Zr are optimally matched, jointly achieving high volume fraction precipitation of nanoscale α2 phase and long-term microstructure stability.

[0010] Mo and V, as β-phase stabilizing elements, introduce an appropriate amount of residual β-phase into the alloy. On the one hand, this buffers the localized stress concentration caused by α2-phase precipitation, improving the alloy's toughness; on the other hand, it enhances the alloy's high-temperature creep resistance through β-phase optimization. The contents of both are controlled at 0.5 wt.%. 2.5 wt.%, ensuring both strength and plasticity.

[0011] Based on the functional division of the aforementioned elements, this invention achieves a systematic integration of performance through the synergistic proportions of five elements: Al, Sn, Zr, Mo, and V. Al strengthens the basic strength framework of the α matrix through solid solution reinforcement; Sn promotes high-density nucleation and precipitation of the α2 phase by increasing the alloy's electron concentration (target value ≤ 2.14), thereby enhancing high-temperature strength; Zr utilizes its infinite solid solubility in Ti to generate lattice distortion and solute dragging effects, reducing the α / α2 phase interface energy and inhibiting atomic diffusion, thus effectively suppressing the coarsening and segregation of the α2 phase during long-term high-temperature heat exposure, ensuring that the α2 phase maintains a uniform nanoscale dispersion distribution before and after heat exposure; Mo and V introduce a small amount of β phase, providing plasticity coordination channels while strengthening the α2 phase, improving the alloy's room-temperature plasticity. The five elements, according to their functional division of solid solution reinforcement (Al), precipitation reinforcement and size stability (Sn+Zr), and plasticity coordination (Mo+V), enable the alloy to achieve a comprehensive match of high strength, high plasticity, and high thermal stability at 600 °C.

[0012] Furthermore, the impurities include C, H, O, and N, with O content ≤ 0.15 wt.%, C content ≤ 0.1 wt.%, N content ≤ 0.05 wt.%, H content ≤ 0.1 wt.%, and the total impurity content not exceeding 0.2 wt.%.

[0013] Furthermore, the near-α type titanium alloy has a high volume fraction of nanoscale α2 phase, which is uniformly dispersed.

[0014] On the other hand, the present invention provides a method for preparing a high-strength, high-thermal-stability near-α-type titanium alloy, the method comprising: S1. Ingot smelting: Using sponge titanium (99.8%), pure aluminum (99.99%), crystalline zirconium (99.95%), Al-V master alloy, Al-Mo master alloy and Ti-Sn master alloy as raw materials, the alloy is smelted under argon atmosphere protection to obtain alloy ingots. The composition of this alloy is also applicable to the alloy ingots obtained by subsequent smelting, that is, the composition of the alloy ingot is the same as the composition of the near-α type titanium alloy. S2, Initial Forging: Remove surface defects from the alloy ingot obtained in S1, and forge it in a muffle furnace at 1000°C. After holding at 1200℃, the billet is forged, with a cumulative deformation of 60%. 70%, after forging and deformation, air-cooled to room temperature to obtain titanium alloy forging billet; S3. Billet Preparation: The titanium alloy forging billet in S2 is deformed by free forging or rotary forging, with a heating temperature of 940°C. 980 ℃, heat preservation time 60 After 180 minutes, the cumulative deformation was 65%. 80%, thus obtaining the near-α type titanium alloy.

[0015] Furthermore, in S1, the V mass fraction of the Al-V master alloy is ≥50 wt.%, the Mo mass fraction of the Al-Mo master alloy is ≥60 wt.%, and the Sn mass fraction of the Ti-Sn master alloy is ≥80 wt.%.

[0016] Furthermore, the melting method in S1 is either cold crucible induction melting or vacuum arc melting.

[0017] Furthermore, the muffle furnace heating in S2 takes place in air, eliminating the need for a high-purity argon protective atmosphere.

[0018] Furthermore, the initial forging in S2 is carried out on a high-speed forging machine or a hydraulic press.

[0019] Furthermore, the S2 forging process employs a five-upsetting and five-drawing process, involving three heating cycles. The first heating cycle includes the first and second upsetting and drawing, with a heating temperature of 1100°C. 1200 ℃, heat preservation time 60 180 min, deformation 30% 50%, deformation time approximately 2 3 min. Second heating: Third and fourth upsetting and drawing, heating temperature 1050℃. 1100 ℃, heat preservation time 30 60 min, deformation 25% 45%, deformation time approximately 2 3 min. Third heating: Fifth upsetting and drawing, heating temperature 1000℃. 1050 ℃, heat preservation time 30 60 min, deformation 15% 35%, deformation time approximately 1 minute 2 min.

[0020] Furthermore, in S3, the final forging temperature of free forging or rotary forging is ≥900 ℃, and after deformation, it is air-cooled to room temperature.

[0021] The beneficial effects of this invention are: 1. Synergistic Strengthening by Dual-State Microstructure and Nanoscale α2 Phase. The near-α type titanium alloy prepared by the method of this invention exhibits a dual-state microstructure, composed of micron-sized equiaxed crystals and nanoscale α2 ordered phase. The nanoscale α2 ordered phase has a high volume fraction and shows a uniform and dispersed distribution in both the forged state and after heat exposure at 600℃ / 100 h. During deformation, the nanoscale α2 phase pins dislocations and strengthens grain boundaries, significantly improving the high-temperature strength of the alloy; simultaneously, its uniform and dispersed distribution and dimensional stability effectively avoid stress concentration and embrittlement, significantly enhancing the thermal stability of the alloy.

[0022] 2. Excellent room temperature and high temperature mechanical properties. The near-α type titanium alloy designed in this invention exhibits excellent room temperature and high temperature mechanical properties. At room temperature, the tensile strength can reach over 1253 MPa, which is more than 19% higher than that of Ti60 alloy, and the elongation after fracture can reach over 18.9%. At 600℃, the tensile strength can reach over 800 MPa, which is more than 30% higher than that of Ti60 alloy, and the elongation after fracture can reach over 38%. After heat exposure at 600℃ for 100 h, the room temperature tensile strength remains almost unchanged, and the elongation after fracture still remains above 14.5%, which is more than 3 / 5 higher than that of Ti60 alloy, achieving a synergistic improvement in high temperature strength and thermal stability.

[0023] 3. The alloy composition is simple, facilitating low-cost manufacturing. The alloy composition of this invention does not contain high-melting-point refractory elements (Nb, Ta, W) or rare earth elements, and does not contain Si. The alloy system is a Ti-Al-Sn-Zr-Mo-V hexa-component system. This avoids metallurgical defects such as segregation and inclusions, as well as increased density problems caused by refractory elements. It also avoids the adverse effects of Si addition on thermal stability and the increased costs associated with rare earth element addition, thus enabling lightweight, low-cost, and high-quality alloy manufacturing.

[0024] 4. The preparation process is simple and suitable for industrial production. The preparation method of near-α type titanium alloy of this invention only requires three main processes: ingot melting, billet forging, and billet preparation. The entire process adopts conventional titanium alloy industrial production equipment (vacuum arc remelting furnace, high-speed forging machine, free forging / rotary forging equipment, etc.), without the need for special processes such as rapid solidification and powder metallurgy. It has a wide process window and strong operability, making it suitable for large-scale industrial production and application of high-temperature titanium alloys.

[0025] 5. Based on the above advantages, this invention abandons the complex alloying route that relies on elements such as Si, Nb, Ta, and W in the prior art. It achieves a better match between strength and plasticity and thermal stability than existing 600℃ high-temperature titanium alloys (such as Ti60) through the synergistic effect of high Sn and high Zr using a simple six-element alloy system. It has good industrial practical value and broad application prospects. Attached Figure Description

[0026] Figure 1 The microstructure of the near-α type titanium alloy prepared in Example 1 in the forged state is shown.

[0027] Figure 2 The microstructure of the near-α type titanium alloy prepared in Example 1 after heat exposure at 600 ℃ for 100 h is shown.

[0028] Figure 3 This is a process flow diagram of the near-α type titanium alloy preparation method of the present invention. Detailed Implementation

[0029] To make the technical problems, technical solutions and advantages of the present invention clearer, a detailed description will be given below in conjunction with the accompanying drawings and specific embodiments.

[0030] Based on the electron concentration theory, this invention designs a novel high-strength, high-thermal-stability near-α-type titanium superalloy characterized by high Sn and high Zr content. Its composition design follows the principles and mechanisms described below: (1) Abandoning the traditional approach of adding Si, instead of relying on silicide strengthening, this design induces α2 phase precipitation by increasing the overall electron concentration of the alloy to achieve high-temperature strengthening. Sn, with its high number of valence electrons (4), is the most effective element for increasing electron concentration. Although Zr has a lower number of valence electrons (2), its infinite solid solubility in α-Ti allows it to reduce the α / α2 phase interface energy through lattice distortion, generating a solute dragging effect that effectively hinders the diffusion of interface atoms, thereby suppressing the coarsening and segregation of the α2 phase. The two complement each other in terms of electron concentration contribution and microstructure stability control, which is the theoretical core of this alloy system design.

[0031] (2) The amount of Al added needs to balance solid solution strengthening effect and plasticity retention. When the Al content is below 5.5 wt.%, the high temperature strength is insufficient; when it is above 7.0 wt.%, the α2 phase is prone to spontaneous precipitation and rapid coarsening during hot working or service, resulting in a sharp decrease in plasticity. Therefore, it is limited to the range of 5.5~7.0 wt.% as the basic strength support.

[0032] (3) The increase in Sn content is the key difference from existing alloys. Existing near-α type titanium alloys at 600 ℃ and 650 ℃ (such as Ti60, Ti600, Ti1100, Ti65, etc.) have a Sn content of 2.7~4.8 wt.% and an electron concentration of about 2.05~2.12. This invention increases the Sn content to 5.5~8.0 wt.%, making the alloy electron concentration break through 2.12 and approach the upper limit of 2.14, thereby significantly improving the thermodynamic driving force for α2 phase precipitation.

[0033] (4) The simultaneous increase of Zr content is a necessary condition for ensuring thermal stability. The Zr content in existing alloys is 2.0~4.0 wt.%, and it usually forms a Ti-Zr-Si compound strengthening phase with Si. This invention increases the Zr content to 5.0~8.0 wt.% without adding Si, and utilizes the solute dragging effect of high-concentration Zr to effectively suppress the coarsening behavior of high volume fraction α2 phase. At the same time, the Zr / Sn ratio is strictly controlled to be 0.8~1.5 to ensure that the precipitation-promoting effect of Sn and the stabilizing effect of Zr are kinetically coordinated and matched, avoiding the instability of the structure due to the excess of one element.

[0034] (5) The amounts of β-stabilizing elements Mo and V added are determined based on the molybdenum equivalent principle, [Mo] 当 =%Mo+%V / 1.4 (wt.%) is controlled within the range of 1.5~2.0. This molybdenum equivalent range can introduce sufficient residual β phase to improve room temperature plasticity without reducing the high-temperature strength of the alloy due to excessive β phase. Mo mainly contributes to high-temperature solid solution strengthening, while V focuses on improving room temperature strength and plasticity, and the two complement each other in the temperature range.

[0035] (6) Electron concentration is the core constraint index for the composition design of this invention. The composition is designed using six elements: Ti, Al, Sn, Mo, V, and Zr, with corresponding valence electron numbers N0 and N1 respectively. Ti = N Zr = 2, N Sn = 4, N Al = N V = 3, N Mo = 5. The electron concentration theory can be expressed by equation (1): (1) in, for The contribution of electron concentration to element atoms, for The atomic percentage of an element in the α phase. When the electron concentration of the alloy is 2.14, the design content range of each alloying element can be determined by combining equation (1), molybdenum equivalent, and Zr / Sn constraint conditions.

[0036] The core difference between the technical approach of this invention and existing near-α type high-temperature titanium alloys lies in the fact that it completely abandons the element Si and no longer relies on the silicide strengthening mechanism. Instead, it significantly increases the content of Sn and Zr at the same time. Sn drives the high-density precipitation of α2 phase to improve high-temperature strength, while Zr inhibits the coarsening of α2 phase to ensure thermal stability. The two work synergistically in terms of electron concentration and kinetic stability, ultimately achieving a uniform and diffuse distribution of high volume fraction nanoscale α2 phase before and after long-term heat exposure, thereby simultaneously improving the high-temperature strength and thermal stability of the alloy.

[0037] Example 1

[0038] A high-strength, high-thermal-stability near-alpha titanium alloy, wherein the mass percentages of each component of the near-alpha titanium alloy are: 6.14 wt.% Al, 5.82 wt.% Zr, 6.15 wt.% Sn, 1.01 wt.% Mo, 1.02 wt.% V, 0.092 wt.% O, 0.007 wt.% C, 0.007 wt.% N, 0.065 wt.% H, with the balance being Ti. The relationship between Zr and Sn is Zr / Sn = 0.95.

[0039] A method for preparing a high-strength, high-thermal-stability near-α-type titanium alloy, the method comprising: S1. Ingot Smelting: Using 7831.81g of sponge titanium (99.8%), 168.33g of Al-60 wt.% Mo master alloy, 768.75g of Ti-80 wt.% Sn master alloy, 582g of crystalline zirconium (99.95%), 290.92g of pure aluminum (99.99%), and 204g of Al-50 wt.% V master alloy as raw materials, a 10 kg alloy ingot was obtained by vacuum consumable melting three times under argon atmosphere protection. S2. Initial Forging: Surface defects are removed from the alloy ingot obtained in S1, and the ingot is subjected to three-stage initial forging using a five-upsetting and five-drawing process on a high-speed forging mill. First Stage: First and second upsetting and drawing, heating temperature 1170 ℃, holding time 60 min, deformation amount 38%, deformation time approximately 3 min. Second Stage: Third and fourth upsetting and drawing, heating temperature 1100 ℃, holding time 30 min, deformation amount 30%, deformation time approximately 3 min. Third Stage: Fifth upsetting and drawing, heating temperature 1050 ℃, holding time 30 min, deformation amount 20%, deformation time approximately 2 min. The total cumulative deformation of the initial forging is 65%. After deformation, the ingot is air-cooled to room temperature to obtain a titanium alloy forging billet.

[0040] S3. Billet preparation: The titanium alloy forging billet in S2 is deformed by free forging. The heating temperature is 950℃, the holding time is 60 min, the cumulative deformation is 70%, and the final forging temperature is about 900℃. After deformation, it is air-cooled to room temperature to obtain the near-α type titanium alloy.

[0041] To compare performance, a Ti60 alloy was prepared using the same method as described in this invention, and its mechanical properties are shown in Table 1. As shown in Table 1, the tensile strength of Example 1 at room temperature and 600 °C was increased by 19% and 30% respectively compared to the Ti60 alloy, and the elongation after fracture was also significantly improved. After heat exposure at 600 °C for 100 h, the room temperature elongation after fracture of Example 1 was increased by more than 3 / 5 compared to the Ti60 alloy. The results indicate that the near-α type titanium alloy prepared in Example 1 of this invention possesses superior thermal strength and thermal stability.

[0042] Table 1. Performance comparison of near-α type titanium alloy and Ti60 alloy prepared in Example 1 of this invention.

[0043] Example 2

[0044] A high-strength, high-thermal-stability near-alpha titanium alloy, wherein the mass percentages of each component of the near-alpha titanium alloy are: 6.46 wt.% Al, 6.95 wt.% Zr, 5.98 wt.% Sn, 1.01 wt.% Mo, 1.02 wt.% V, 0.097 wt.% O, 0.089 wt.% C, 0.009 wt.% N, 0.007 wt.% H, with the balance being Ti. The relationship between Zr and Sn is Zr / Sn = 1.31.

[0045] A method for preparing a high-strength, high-thermal-stability near-α-type titanium alloy, the method comprising: S1. Ingot Smelting: Using 7703.71g sponge titanium (99.8%), 168.33g Al-60 wt.% Mo master alloy, 747.5g Ti-80 wt.% Sn master alloy, 695.35g crystalline zirconium (99.95%), 476.71g pure aluminum (99.99%) and 204g Al-50 wt.% V master alloy as raw materials, a 10 kg alloy ingot was obtained by vacuum consumable melting three times under argon atmosphere protection. S2. Initial Forging: Surface defects were removed from the alloy ingot obtained in S1, and the ingot was subjected to three-stage initial forging using a five-upsetting and five-drawing process on a high-speed forging mill. First Stage: First and second upsetting and drawing, heating temperature 1130 ℃, holding time 60 min, deformation amount 39%, deformation time approximately 3 min. Second Stage: Third and fourth upsetting and drawing, heating temperature 1070 ℃, holding time 30 min, deformation amount 31%, deformation time approximately 3 min. Third Stage: Fifth upsetting and drawing, heating temperature 1010 ℃, holding time 30 min, deformation amount 21%, deformation time approximately 2 min. The total cumulative deformation was 68%. After deformation, the ingot was air-cooled to room temperature to obtain a titanium alloy forging billet. S3. Billet preparation: The titanium alloy forging billet in S2 is deformed by rotary forging. The heating temperature is 970 ℃, the holding time is 120 min, the cumulative deformation is 80%, and the final forging temperature is about 920 ℃. After deformation, it is air-cooled to room temperature to obtain the near-α type titanium alloy.

[0046] Table 2. Performance comparison of near-α type titanium alloy and Ti60 alloy prepared in Example 2 of this invention.

[0047] Comparative Example In order to objectively evaluate the advancement of the technical solution of the present invention compared with the prior art, this comparative example selects the composition of a typical near-α type high-temperature titanium alloy in the prior art and prepares the comparative alloy using the same preparation method as the embodiment of the present invention.

[0048] The composition of a typical near-α type high-temperature titanium alloy in the prior art is as follows: Al 5.72 wt.%, Sn 4.01 wt.%, Zr 3.93 wt.%, Mo 0.25 wt.%, Si 0.39 wt.%, Nb 0.41 wt.%, Ta 2.54 wt.%, W 0.053 wt.%, Fe 0.03 wt.%, O 0.08 wt.%, with the balance being Ti. The same preparation method as in Example 1 of this invention was used: ingot melting → billet forging → billet preparation to prepare the comparative alloy. The billet forging temperature was 1000℃~1200℃, the billet preparation heating temperature was 940℃~980℃, and the final forging temperature was ≥900℃.

[0049] Table 3 shows a comparison of the mechanical properties of the alloy with those of Examples 1 and 2 of this invention.

[0050] Table 3 Performance Comparison of Embodiments and Comparative Examples of the Invention

[0051] As shown in Table 3, the room temperature and 600°C tensile strengths of Examples 1 and 2 of the present invention are significantly superior to those of the comparative alloy. The room temperature elongation after fracture is increased by more than 45% and 12% respectively compared to the comparative alloy, and the 600°C elongation after fracture is increased by more than 73% respectively. Particularly important, after heat exposure at 600°C for 100 hours, the room temperature elongation after fracture of Examples 1 and 2 of the present invention reaches 14.5% and 11.9% respectively, while the comparative alloy only retains 6.5%. The room temperature elongation after fracture of the alloy of the present invention after heat exposure is increased by more than 123% and 83% respectively compared to the comparative alloy.

[0052] This indicates that, by abandoning the addition of elements such as Si, Nb, Ta, and W in traditional alloys, the present invention, through the synergistic effect of high Sn and high Zr, not only significantly improves the room temperature and high temperature strength of the alloy, but also achieves a breakthrough improvement in thermal stability, achieving a strength-plasticity match superior to existing typical near-α type high temperature titanium alloys.

[0053] The above description represents the preferred embodiments of the present invention. It should be noted that those skilled in the art can make various improvements and modifications without departing from the principles of the present invention, and these improvements and modifications should also be considered within the scope of protection of the present invention.

Claims

1. A high-strength, high-thermal-stability near-α-type titanium alloy, characterized in that, The near-alpha titanium alloy is composed of the following components by mass percentage: Al 5.5 wt.% 7.0 wt.%, Sn 5.5 wt.% 8.0 wt.%, Zr 5.0 wt.% 8.0 wt.%, Mo 0.5 wt.% 2.5 wt.%, V 0.5 wt.% 2.5 wt.%, the balance being Ti and unavoidable impurities; and the mass ratio of Zr to Sn satisfies 0.8≤Zr / Sn≤1.

5.

2. The high-strength, high-thermal-stability near-α-type titanium alloy according to claim 1, characterized in that, The unavoidable impurities include C, H, O, and N, wherein the O content is ≤0.15 wt.%, the C content is ≤0.1 wt.%, the N content is ≤0.05 wt.%, the H content is ≤0.1 wt.%, and the total impurity content does not exceed 0.2 wt.%.

3. The high-strength, high-thermal-stability near-α-type titanium alloy according to claim 1, characterized in that, The near-α type titanium alloy contains a high volume fraction of nanoscale α2 ordered phase.

4. The high-strength, high-thermal-stability near-α-type titanium alloy according to claim 3, characterized in that, The high volume fraction of nanoscale α2 ordered phase exhibits a uniformly dispersed distribution in both the forged state and after heat exposure at 600℃ / 100h.

5. The high-strength, high-thermal-stability near-α-type titanium alloy according to claim 1, characterized in that, The mass ratio of Zr to Sn satisfies 0.8 ≤ Zr / Sn ≤ 1.

5.

6. The high-strength, high-thermal-stability near-α-type titanium alloy according to claim 1, characterized in that, The near-α type titanium alloy is a forged alloy.

7. A method for preparing a high-strength, high-thermal-stability near-α-type titanium alloy according to any one of claims 1 to 6, characterized in that, Includes the following steps: S1. A titanium source, an aluminum source, a zirconium source, a vanadium source, a molybdenum source, and a tin source are smelted under a protective atmosphere to obtain an alloy ingot; the composition of the alloy ingot is as described in claim 1. S2. The alloy ingot obtained in S1 is heated at 1000℃. Forging was carried out at 1200℃, with a cumulative deformation of 60%. 70%, after forging and deformation, air-cooled to room temperature to obtain titanium alloy forging billet; S3, heat the titanium alloy forging billet from S2 at 940℃ Deformation processing was performed at 980℃, with a holding time of 60 min. After 180 minutes, the cumulative deformation was 65%. 80%, to obtain the near-α type titanium alloy.

8. The preparation method according to claim 7, characterized in that, The titanium source in S1 is sponge titanium, the aluminum source is pure aluminum, the zirconium source is crystalline zirconium, the vanadium source is an Al-V master alloy, the molybdenum source is an Al-Mo master alloy, and the tin source is a Ti-Sn master alloy.

9. The preparation method according to claim 7, characterized in that, The melting method in S1 is vacuum self-consumable melting, cold crucible induction melting, or vacuum arc melting; the protective atmosphere is an inert gas atmosphere.

10. The preparation method according to claim 7, characterized in that, The forging process described in S2 employs multi-fire upsetting and drawing forging, wherein the first fire involves the first and second upsetting and drawing processes, with a heating temperature of 1100°C. 1200 ℃, heat preservation time 60 180 min, deformation 30% 50%, deformation time approximately 2 3 min. Second heating: Third and fourth upsetting and drawing, heating temperature 1050℃. 1100 ℃, heat preservation time 30 60 min, deformation 25% 45%, deformation time approximately 2 3 min. Third heating: Fifth upsetting and drawing, heating temperature 1000℃. 1050 ℃, heat preservation time 30 60 min, deformation 15% 35%, deformation time approximately 1 minute 2 min.

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