Ultrahigh-strength titanium alloy with TRIP effect and preparation method thereof

Abstract

The ultrahigh-strength titanium alloy with the TRIP effect comprises the following chemical components in percentage by weight: 5.0 to 7.0 percent of Al; 3.0 to 5.0 parts of V; 3.0 to 7.0 parts of Cu; 0.3 to 0.7 part of Si; 0.15 to 0.35 part of O; and the balance of Ti. The preparation method of the ultrahigh-strength titanium alloy with the TRIP effect comprises the following steps: smelting the raw materials according to the formula in a consumable electrode vacuum furnace, and then forging the raw materials into a blank at the temperature of 1000 DEG C or above; the forging blank is subjected to heat preservation for 1-3 hours at the temperature of 1000 DEG C or above and then quenched to obtain a full-martensite structure; the blank is subjected to hot rolling in a (alpha + beta + Ti5Si3) three-phase region at the temperature of 770-850 DEG C, and a nanometer ultra-fine grain structure is obtained; and after solid solution at 800-840 DEG C, quenching is performed, so that a Ti2Cu phase separated out in the rolling and cooling process is redissolved in a matrix, and the TRIP effect is given to the material. The method is high in dynamic recrystallization nucleation rate and has an ultra-fine grain structure; the proportion and fault energy of the beta phase can be regulated and controlled, and the finished product has an excellent TRIP effect.

Description

Technical Field

[0001] The present invention relates to the application technical fields of titanium alloy material scheme design and preparation, and in particular to an ultra-high-strength titanium alloy with TRIP effect and a preparation method thereof. Background Art

[0002] Ultra-high-strength titanium alloys with the TRIP effect are titanium alloys that significantly enhance both strength and ductility through the transformation-induced plasticity (TRIP) mechanism. The TRIP effect primarily relies on the stress-induced martensitic transformation in β-phase titanium alloys. This transformation significantly increases the material's strain hardening rate and uniform elongation, achieving a combination of high strength and high ductility.

[0003] The application of the TRIP effect in titanium alloys has been extensively researched and developed. For example, a research team at the Beijing University of Aeronautics and Astronautics has developed a new TRIP / TWIP titanium alloy with a strain hardening rate of up to 6.1 GPa, which provides new ideas for designing high-strength and high-ductility titanium alloys. Furthermore, by controlling the alloy's β stability and adding elements such as Zr, the TRIP effect can be further optimized, thereby improving the alloy's yield strength and work hardening ability.

[0004] Although the TRIP effect can significantly enhance the performance of titanium alloys, its practical application still faces some challenges. For example, the TRIP effect is often accompanied by a decrease in yield strength, which limits its application in certain high-strength scenarios. Therefore, how to enhance yield strength while maintaining high strain hardening capacity is an important direction for future research.

[0005] Ultra-high-strength titanium alloys with the TRIP effect significantly enhance their strength and ductility through a phase transformation-induced plasticity mechanism. However, practical applications still require a balance between yield strength and strain hardening. Future research could further enhance the comprehensive performance of these alloys by optimizing alloy composition, manipulating microstructure, and integrating multi-scale deformation mechanisms.

[0006] Several strategies can be adopted to enhance the yield strength of titanium alloys with TRIP effect by optimizing alloy composition: Strategy 1, adding alloying elements: For example, adding titanium, niobium, and vanadium, these elements can increase the strength of titanium alloys through precipitation hardening mechanisms while having little effect on weldability and formability. In 3D-printed Ti-5553 alloy, the addition of Mo nanoparticles significantly improved the mechanical properties and uniformity, thereby increasing the yield strength.

[0007] Strategy 2: Controlling grain size: Reducing the grain size can significantly increase the yield strength of titanium alloys. For example, in the metastable β-type Ti-4V-2Mo-2Fe alloy, when the grain size is 13μm, the yield strength can exceed 900MPa.

[0008] Strategy 3, heterogeneous laminate structure design: By regulating the thickness of the matrix layer and the distribution of the hard-plastic nano-precipitated phase, the activation order of TRIP and ODP can be precisely controlled. ODP can be activated first at low strain and then TRIP can be activated at high strain, thereby achieving the synergistic coupling of the deformation-phase transformation mechanism of metastable titanium alloys, doubling the yield strength of the alloy without losing its uniform plasticity.

[0009] Strategy 4, natural aging: In Ti-15Nb-5Zr-4Sn-1Fe alloy, natural aging leads to a significant increase in yield strength from 493 MPa to 683 MPa without sacrificing ductility. Natural aging reduces the reversibility of ω to β through the diffusion and coarsening of ω particles and hinders the stress-induced β to α' martensitic phase transformation under subsequent deformation.

[0010] Strategy 5, cold rolling: Cold rolling of the commercial β-Ti alloy Ti-10V-2Fe-3Al can significantly increase its yield strength while maintaining high strain hardenability and uniform tensile ductility. For example, a 5% cold rolling condition provides the best yield strength (about 900 MPa), ductility (8%), and strain hardening.

[0011] People are eager to obtain an ultra-high-strength titanium alloy with TRIP effect and a preparation method thereof with better technical effects. Summary of the Invention

[0012] The present invention provides an ultra-high-strength titanium alloy with TRIP effect and a preparation method thereof.

[0013] An ultra-high-strength titanium alloy with TRIP effect, the key technology is: Calculated by weight percentage, the chemical composition and content of the titanium alloy meet the following requirements: Al: 5.0-7.0; V: 3.0-5.0; Cu: 3.0-7.0; Si: 0.3-0.7; O: 0.15-0.35; and the balance is Ti.

[0014] Further preferred technical requirements are: calculated by weight percentage, the chemical composition and content of the titanium alloy meet the following requirements: Al: 5.8~6.3; V: 3.7~4.2; Cu: 4.5~5.5; Si: 0.45~0.55; O: 0.22~0.27; the balance is Ti; the content of impurity elements in the alloy complies with the corresponding requirements in the national standard GB / T 3620.1-2016 "Table of Titanium and Titanium Alloy Grades and Chemical Compositions".

[0015] Preferably, the grade of the ultra-high strength titanium alloy with TRIP effect is Ti6Al4V5Cu0.25Si0.25O.

[0016] The present invention also claims protection for a method for preparing the ultra-high-strength titanium alloy with TRIP effect, the key technical aspects of which are: Step 1: using a vacuum consumable furnace to melt an ingot: the raw materials used for melting the ingot are as follows according to the chemical composition and weight percentage: Al: 5.8-6.3; Cu: 4.5-5.5; V: 3.7-4.2; Si: 0.45-0.55; O: 0.22-0.27; the balance is Ti; It is then forged into billets at temperatures above 1000°C; Step 2: keeping the forging blank at above 1000° C. for 1 to 3 hours and then quenching to obtain a fully martensitic structure; Step 3: by adding Cu, Si, and O elements, during the quenching process, Cu, Si, and O are supersaturated and dissolved in the β matrix, causing strong lattice distortion and refining the width of the martensite laths; the width of the martensite laths after quenching is refined to obtain an ultrafine lath martensite structure; Step 4, using the aforementioned ultrafine lath martensite as the initial structure, hot rolling the billet in the (α+β+Ti5Si3) three-phase region at 770-850°C to obtain a nano-ultrafine grain structure; Step 5: The rolled and deformed material is subjected to solution treatment at 800-840°C and then quenched, so that the Ti2Cu phase precipitated during the rolling cooling process dissolves back into the matrix, giving the material a TRIP effect.

[0017] The method for preparing the ultra-high-strength titanium alloy with TRIP effect of the present invention preferably claims the following technical contents: During the rolling process of step 4, the deformation amount is required to be no less than 70%; further preferably, during the rolling process of step 4, the rolling deformation amount is required to be no less than 80%.

[0018] The rolling deformation temperature is 800~820℃, and the solution treatment temperature after rolling is 810~820℃.

[0019] Different from the prior art, the titanium alloy provided by the present invention does not need to rely on high-power equipment and expensive molds, and can achieve the preparation of ultrafine-grained titanium alloy through conventional thermal deformation.

[0020] The present invention significantly improves the strength of titanium alloy materials without reducing the plastic toughness of the material. The microstructure of the ultra-high-strength titanium alloy with TRIP effect is an ultrafine-grained structure with an α grain size of 230-350 nm, and a metastable β phase can be observed around the α grain boundaries. Its room temperature yield strength is 720-900 MPa, the tensile strength is 1150-1620 MPa, the yield strength ratio is 0.55-0.65, the elongation is 7-25%, and the U-mouth impact energy is 15-70 J / cm 2 .

[0021] To reconcile the inherent contradiction between high strength and high plasticity and toughness in titanium alloys and develop a new titanium alloy with excellent comprehensive mechanical properties, we designed and prepared an ultrafine-grained Ti6Al4V5Cu0.5Si0.25O titanium alloy with the TRIP effect. This ultrafine-grained structure restricts dislocation motion, enhancing the material's strength. Furthermore, the metastable β phase distributed along grain boundaries generates a transformation-induced plasticity (TRIP) effect under loading, further enhancing the material's plasticity and toughness.

[0022] The present invention adds Cu, Si, and O to a conventional TC4 titanium alloy, dissolving them in the β single-phase region. Through quenching and rapid cooling, the Cu, Si, and O elements are supersaturated in the matrix, inducing strong lattice distortion. This adaptive lattice distortion ultimately results in an ultrafine lath martensite structure after quenching. This ultrafine lath martensite structure is then hot-rolled in the (α+β+Ti5Si3) three-phase region at 770-850°C. The strain energy and interfacial energy stored in the martensite provide a strong driving force for dynamic recrystallization of the α phase, significantly increasing the dynamic recrystallization nucleation rate. Simultaneously, during the hot rolling process, the martensite decomposes from α′ to α+β+Ti5Si3. The decomposition product, β / Ti5Si3, pins the α grains, inhibiting the coarsening and growth of dynamic recrystallization nuclei, resulting in an ultrafine-grained structure after deformation.

[0023] To allow the Ti2Cu phase precipitated during the slow cooling process after rolling to dissolve back into the matrix, the present invention subjects the rolled titanium alloy to a solution-hardening treatment followed by quenching. This treatment modulates the β-phase ratio and stacking fault energy, ultimately imparting the TRIP effect to the material, resulting in the preparation of an ultra-high-strength Ti6Al4V5Cu0.5Si0.25O titanium alloy exhibiting the TRIP effect. BRIEF DESCRIPTION OF THE DRAWINGS

[0024] Figure 1 This is a SEM photograph of the titanium alloy forging blank after quenching in Example 5; Figure 2 This is an SEM photograph of the titanium alloy forging blank after quenching in Comparative Example 10; Figure 3 This is a SEM photograph of the solution quenched structure of the titanium alloy in Example 5 after hot rolling; Figure 4 This is a SEM photograph of the titanium alloy after hot rolling in Comparative Example 10; Figure 5 This is a TEM high-resolution observation photograph of the structure of the titanium alloy in Example 5 after room temperature tensile testing. DETAILED DESCRIPTION

[0025] The present invention will be further described below with reference to the embodiments and the accompanying drawings, but is not limited thereto.

[0026] The present invention provides a novel titanium alloy having the following chemical composition: Al: 5.0-7.0; V: 3.0-5.0; Cu: 3.0-7.0; Si: 0.3-0.7; O: 0.15-0.35; the balance being Ti. The content of impurity elements in the alloy must comply with the requirements of the national standard GB / T 3620.1-2016, "Titanium and Titanium Alloy Grades and Chemical Composition Table."

[0027] See also Figures 1 and 2 , Figure 1 is a SEM photograph of the titanium alloy forging blank after quenching in Example 5 of the present invention, Figure 2 This is an SEM photograph of the titanium alloy forged billet after quenching in Comparative Example 10 of the present invention. It can be seen that by adding Cu, Si, and O elements to the traditional TC4 titanium alloy, the size of the martensitic laths after forging and quenching can be reduced, resulting in an ultrafine nano-lath martensitic structure.

[0028] According to the method provided by this invention, after quenching and hot rolling of titanium alloy, a nano-ultrafine grain structure is formed, and the β phase, Ti5Si3 phase, and Ti2Cu phase are distributed along the grain boundaries after hot rolling, which can effectively inhibit the coarsening and growth of α grains during rolling deformation.

[0029] Figure 3 This is a SEM photograph of the solution quenched structure of the titanium alloy in Example 5 of the present invention after hot rolling. It can be seen that the Ti2Cu phase precipitated during the rolling process has dissolved back into the matrix after solution quenching.

[0030] Figure 4 This is a SEM photograph of the hot-rolled titanium alloy in Comparative Example 10. It can be seen that the α grain size of the traditional TC4 titanium alloy after hot rolling is several microns, and the size of the β phase is also significantly larger than that of the titanium alloy provided by the present invention.

[0031] Figure 5TEM observation of the β phase around the grain boundary of the titanium alloy in Example 5 of the present invention after tensile deformation shows that the β phase transforms into the α′′ phase with martensite characteristics after stretching, indicating that the material has a transformation induced plasticity (TRIP) effect.

[0032] The present application will be described and explained below through several groups of specific embodiments and comparative examples, but they should not be used to limit the scope of the present application.

[0033] Examples: Examples 1 to 9 are Ti6Al4V5Cu0.5Si0.25O alloys smelted according to the chemical composition range provided by the present invention, wherein the contents of Al, V, Cu, Si, and O elements are gradually increased, and the corresponding preparation processes are appropriately adjusted within the technical parameter range specified in the present invention.

[0034] Comparative Examples: The chemical composition of Comparative Example 1 is lower than the lower limit of the chemical composition range provided by the present invention, while the chemical composition of Comparative Example 9 is higher than the upper limit of the chemical composition range provided by the present invention. By comparing them with Examples 1 and 9, the effects of chemical composition on the microstructure and properties of titanium alloys are illustrated. In Comparative Example 2, the quenching temperature of the forging blank is lower than the temperature specified in the present invention. By comparing it with Example 2, the effects of the quenching temperature of the forging blank on the microstructure and properties of titanium alloys are illustrated. In Comparative Example 3, the cooling method after solution treatment of the forging blank is not the quenching specified in the present invention, but furnace cooling. By comparing it with Example 3, the effects of the cooling method after solution treatment of the forging blank on the microstructure and properties of titanium alloys are illustrated. In Comparative Example 4, the rolling deformation temperature is higher than the upper limit specified in the present invention, while the rolling deformation temperature of Comparative Example 5 is lower than the lower limit specified in the present invention. By comparing them with Examples 4 and 5, the effects of rolling temperature on the microstructure and properties of titanium alloys are illustrated. In Comparative Example 6, the rolling deformation amount is lower than the lower limit specified in the present invention. By comparing it with Example 6, the effects of deformation amount on the microstructure and properties of titanium alloys are illustrated. The quenching temperature of the material in Comparative Example 7 during the solution treatment after rolling was below the lower limit specified in the present invention. The quenching temperature of the material in Comparative Example 8 during the solution treatment after rolling was above the upper limit specified in the present invention. By comparing these with Examples 7 and 8, the effect of the solution treatment temperature after rolling on the microstructure and properties of the titanium alloy is demonstrated. Comparative Example 10 is a commercial TC4 titanium alloy. By comparing it with Examples 1 to 9 of the present invention, the advantages of the microstructure and properties of the titanium alloy provided by the present invention are demonstrated.

[0035] Table 1 Chemical composition and preparation process of materials in Examples and Comparative Examples

[0036] The preparation processes of Examples 1-9 and Comparative Examples 1-10 are described as follows: Example 1: An ingot was forged at 1000°C, then held at 1000°C for 1 hour and water-quenched to room temperature to produce a martensitic billet. The martensitic billet was hot-rolled at 770°C, with a rolling deformation greater than 70%. The deformed material was water-quenched at 800°C.

[0037] Example 2: The ingot was forged at 1000°C, then held at 1000°C for 1 hour and water-quenched to room temperature to produce a martensitic billet. The martensitic billet was hot-rolled at 780°C, with a rolling deformation greater than 70%. The deformed material was water-quenched at 800°C.

[0038] Example 3: The ingot was forged at 1050°C, then held at 1050°C for 2 hours and water quenched to room temperature to produce a martensitic billet. The martensitic billet was hot rolled at 790°C with a rolling deformation greater than 70%. The deformed material was water quenched at 810°C.

[0039] Example 4: The ingot was forged at 1050°C, then held at 1050°C for 2 hours and water quenched to room temperature to produce a martensitic billet. The martensitic billet was hot rolled at 800°C, with a rolling deformation greater than 80%. The deformed material was water quenched at 810°C.

[0040] Example 5: The ingot was forged at 1100°C, then held at 1100°C for 2 hours and water quenched to room temperature to produce a martensitic billet. The martensitic billet was hot rolled at 810°C, with a rolling deformation greater than 80%. The deformed material was water quenched at 820°C.

[0041] Example 6: The ingot was forged at 1100°C, then held at 1100°C for 2 hours and oil-quenched to room temperature to produce a martensitic billet. The martensitic billet was hot-rolled at 820°C, with a rolling deformation greater than 80%. The deformed material was water-quenched at 820°C.

[0042] Example 7: The ingot was forged at 1150°C, then held at 1150°C for 2 hours and oil-quenched to room temperature to produce a martensitic billet. The martensitic billet was hot-rolled at 830°C, with a rolling deformation greater than 80%. The deformed material was water-quenched at 830°C.

[0043] Example 8: The ingot was forged at 1150°C, then held at 1150°C for 2 hours and oil-quenched to room temperature to produce a martensitic billet. The martensitic billet was hot-rolled at 840°C, with a rolling deformation greater than 85%. The deformed material was water-quenched at 830°C to impart the TRIP effect.

[0044] Example 9: The ingot was forged at 1150°C, then held at 1150°C for 3 hours and oil-quenched to room temperature to produce a martensitic billet. The martensitic billet was hot-rolled at 850°C to a deformation greater than 85%. The deformed material was water-quenched at 840°C.

[0045] Comparative Example 1: An ingot was forged at 1000°C, then held at 1000°C for 1 hour and water-quenched to room temperature to produce a martensitic billet. The martensitic billet was hot-rolled at 770°C, with a rolling deformation greater than 70%. The deformed material was water-quenched at 800°C.

[0046] Comparative Example 2: The ingot was forged at 1000°C, then held at 800°C for 1 hour and water quenched to room temperature to produce a martensitic billet. The martensitic billet was hot rolled at 780°C, with a rolling deformation greater than 70%. The deformed material was water quenched at 800°C.

[0047] Comparative Example 3: The ingot was forged at 1050°C, then held at 1050°C for 2 hours and then slowly cooled to room temperature. The ingot was hot rolled at 790°C, with a rolling deformation greater than 70%. The deformed material was water quenched at 810°C.

[0048] Comparative Example 4: The ingot was forged at 1050°C, then held at 1050°C for 2 hours and water quenched to room temperature to produce a martensitic billet. The martensitic billet was hot rolled at 680°C, with a rolling deformation greater than 80%. The deformed material was water quenched at 810°C.

[0049] Comparative Example 5: The ingot was forged at 1100°C, then held at 1100°C for 2 hours and water quenched to room temperature to produce a martensitic billet. The martensitic billet was hot rolled at 950°C, with a rolling deformation greater than 80%. The deformed material was water quenched at 820°C.

[0050] Comparative Example 6: The ingot was forged at 1100°C, then held at 1100°C for 2 hours and oil-quenched to room temperature to produce a martensitic billet. The martensitic billet was hot-rolled at 820°C with a rolling deformation of 40%. The deformed material was then water-quenched at 820°C.

[0051] Comparative Example 7: The ingot was forged at 1150°C, then held at 1150°C for 2 hours and oil-quenched to room temperature to produce a martensitic billet. The martensitic billet was hot-rolled at 830°C, with a rolling deformation greater than 80%. The deformed material was water-quenched at 500°C.

[0052] Comparative Example 8: The ingot was forged at 1150°C, then held at 1150°C for 2 hours and oil quenched to room temperature to produce a martensitic billet. The martensitic billet was hot rolled at 840°C, with a rolling deformation greater than 85%. The deformed material was water quenched at 950°C.

[0053] Comparative Example 9: The ingot was forged at 1150°C, then held at 1150°C for 3 hours and oil-quenched to room temperature to produce a martensitic billet. The martensitic billet was hot-rolled at 850°C, with a rolling deformation greater than 85%. The deformed material was water-quenched at 840°C.

[0054] Comparative Example 10: The ingot was forged at 1100°C, then kept at 1100°C for 2 hours and then water quenched to room temperature to obtain a martensitic billet. The martensitic billet was hot rolled at 900°C with a rolling deformation greater than 80%.

[0055] Related content description: 1. Microstructure characterization: The microstructure of the material was characterized using a TESCAN scanning electron microscope (SEM). After mechanical grinding and polishing, the SEM test specimen was etched with Kroll etching solution for 20 seconds to observe its microstructure morphology. The experimental results are shown in the figure. Figures 1 and 2 , Figure 5 The microstructure of the material was characterized using a Talos F200x transmission electron microscope (TEM). After the sample was manually thinned to 50 nm, it was chemically double-sprayed in a 10 vol.% HClO4 + 90 vol.% C2H5OH solution. The voltage during thinning was set at 25 V and the temperature was -25 ° C. The experimental results are shown in the figure. Figures 3 and 4 shown.

[0056] 2. Mechanical Properties Testing: φ5M10 standard cylindrical tensile specimens were processed in accordance with GBT 228.1-2010. The room temperature tensile properties of the material were then tested using an Instron 8872 tensile testing machine. The crosshead displacement rate of the tensile test was 0.5 mm / min. The average of three measurements was taken as the tensile properties of the material. The tensile strength, yield strength, and elongation are shown in Table 2. U-notch impact specimens with dimensions of 10 × 10 × 55 mm were processed in accordance with GB / T 229-2020. The impact toughness of the material was tested using a HIT450P impact testing machine. The final test results are given as the average of three measurements, as shown in Table 2.

[0057] From the results in Table 2, it can be seen that Examples 1 to 9 all have α ultrafine grains + metastable β phase structures distributed along the grain boundaries. The materials in the examples all have high strength and good plasticity and toughness. Their yield strength ratio is only 0.55 to 0.65, indicating that they all have TRIP effect. With the increase of Al, V, Cu, Si, and O element content, the tensile strength σ b Significantly improved, while the elongation δ and impact toughness α k Gradually decrease.

[0058] The Al, V, Cu, Si, and O contents of Comparative Example 1 are all below the lower limits specified in the present invention, resulting in relatively coarse grains, a high yield ratio, and no TRIP effect. The material's strength and ductility are significantly lower than those of Example 1 of the present invention. The Al, V, Cu, Si, and O contents of Comparative Example 9 are all above the upper limits specified in the present invention. Although its ultrafine grain structure imparts high strength to the material, its ductility and ductility are significantly lower than those of Comparative Example 9 of the present invention.

[0059] The quenching temperature of the forging blank in Comparative Example 2 is lower than the temperature specified in the present invention, and its structure is a dual-state structure, rather than the α ultrafine grains + metastable β phase structure distributed along the grain boundaries in Example 2 of the present invention. Therefore, its strength and plastic toughness are significantly lower than those in Example 2 of the present invention.

[0060] In Comparative Example 3, the cooling method after solution treatment of the forged blank was not quenching as specified in the present invention, but rather furnace cooling. This resulted in a coarse grain size in the prepared material and the lack of TRIP effect. Consequently, its strength and ductility were significantly lower than those of Example 3 of the present invention.

[0061] Table 2 Microstructure characteristics and mechanical properties of materials in Examples and Comparative Examples

[0062] The rolling deformation temperature of Comparative Example 4 is higher than the upper limit specified in the present invention, and the rolling deformation temperature of Comparative Example 5 is lower than the lower limit specified in the present invention. They both obtain a basketweave structure after rolling deformation, so their strength and plastic toughness are significantly lower than those of Examples 4 and 5 of the present invention.

[0063] The rolling deformation of Comparative Example 6 is lower than the lower limit of deformation specified in the present invention. Therefore, a lamellar structure is obtained after deformation, rather than the α ultrafine grains + metastable β phase structure distributed along the grain boundaries of the present invention. Therefore, its strength and plastic toughness are significantly lower than those of Example 6 of the present invention.

[0064] The material of Comparative Example 7, during solution treatment after rolling, was quenched at a temperature below the lower limit specified in the present invention. Its microstructure is a core-shell structure, rather than the present invention's structure of ultrafine α grains and metastable β phase distributed along grain boundaries. Therefore, while its strength is higher, its plasticity and toughness are significantly lower than those of Example 7 of the present invention.

[0065] When the material of Comparative Example 8 is solution treated after rolling, its quenching temperature is higher than the upper limit specified in the present invention, and its structure is martensite structure, rather than the α ultrafine grain + metastable β phase structure distributed along the grain boundaries of the present invention. Therefore, although its strength is higher, its plastic toughness is significantly lower than that of Example 8 of the present invention.

[0066] Comparative Example 10 is a commercial TC4 titanium alloy having an ultrafine grain structure. By comparing with Examples 1 to 9 of the present invention, it can be seen that the comprehensive mechanical properties of the material provided by the present invention are significantly better than those of the commercial ultrafine grain TC4 titanium alloy.

[0067] Example 10 An ultra-high-strength titanium alloy with TRIP effect, the key technology is: The chemical composition and weight percentage content of the titanium alloy meet the following requirements by weight percentage: Al: 5.8-6.3; V: 3.7-4.2; Cu: 4.5-5.5; Si: 0.45-0.55; O: 0.22-0.27; the balance is Ti; The content of impurity elements in the alloy meets the corresponding requirements of the national standard GB / T3620.1-2016 "Titanium and Titanium Alloy Grades and Chemical Composition Table".

[0068] Preferably, the grade of the ultra-high strength titanium alloy with TRIP effect is Ti6Al4V5Cu0.25Si0.25O.

[0069] Example 11 The preparation method of ultra-high strength titanium alloy with TRIP effect has the following steps and contents: Step 1: using a vacuum consumable furnace to melt the ingot: the raw materials used for melting the ingot are as follows according to the chemical composition and weight percentage: Al: 5.8-6.3; V: 3.7-4.2; Cu: 4.5-5.5; Si: 0.45-0.55; O: 0.22-0.27; the balance is Ti; It is then forged into billets at temperatures above 1000°C; Step 2: keeping the forging blank at above 1000° C. for 1 to 3 hours and then quenching to obtain a fully martensitic structure; Step 3: By adding Cu, Si, and O elements, the width of the martensite lath after quenching can be refined to obtain an ultrafine lath martensite structure; Step 4, using the aforementioned ultrafine lath martensite as the initial structure, hot rolling the billet in the (α+β+Ti5Si3) three-phase region at 770-850°C to obtain a nano-ultrafine grain structure; Step 5: The rolled and deformed material is subjected to solution treatment at 800-840°C and then quenched, so that the Ti2Cu phase precipitated during the rolling cooling process dissolves back into the matrix, giving the material a TRIP effect.

[0070] During the rolling process of step 4, the rolling deformation is required to be no less than 80%.

[0071] The rolling deformation temperature is 800~820℃, and the solution treatment temperature after rolling is 810~820℃.

[0072] Different from the prior art, the titanium alloy provided in this embodiment does not require high-power equipment and expensive molds, and can be prepared by conventional thermal deformation to achieve ultrafine-grained titanium alloy.

[0073] This embodiment significantly improves the strength of titanium alloy materials without reducing the plastic toughness of the material. The microstructure of the ultra-high-strength titanium alloy with TRIP effect is an ultrafine-grained structure with an α grain size of 230-350 nm, and a metastable β phase can be observed around the α grain boundaries. Its room temperature yield strength is 720-900 MPa, the tensile strength is 1150-1620 MPa, the yield strength ratio is 0.55-0.65, the elongation is 7-25%, and the U-mouth impact energy is 15-70 J / cm 2 .

[0074] To reconcile the inherent contradiction between high strength and high plasticity and toughness in titanium alloys and develop a new titanium alloy with excellent comprehensive mechanical properties, this example designed and prepared an ultrafine-grained Ti6Al4V5Cu0.5Si0.25O titanium alloy exhibiting the TRIP effect. Its ultrafine grain structure restricts dislocation motion, enhancing the material's strength. Furthermore, the metastable β phase distributed along grain boundaries generates a transformation-induced plasticity (TRIP) effect under loading, enhancing the material's plasticity and toughness.

[0075] This example adds Cu, Si, and O to a conventional TC4 titanium alloy, dissolving them in the β single-phase region. Through quenching and rapid cooling, the Cu, Si, and O elements are supersaturated in the matrix, inducing strong lattice distortion. This adaptive lattice distortion ultimately results in an ultrafine lath martensite structure after quenching. This ultrafine lath martensite structure is then hot-rolled in the (α+β+Ti5Si3) three-phase region at 770-850°C. The strain energy and interfacial energy stored in the martensite provide a strong driving force for dynamic recrystallization of the α phase, significantly increasing the dynamic recrystallization nucleation rate. Simultaneously, during hot rolling, the martensite decomposes from α′ to α+β+Ti5Si3. The decomposition product, β / Ti5Si3, pins the α grains, inhibiting the coarsening and growth of dynamic recrystallization nuclei, resulting in an ultrafine-grained structure after deformation.

[0076] To allow the Ti2Cu phase precipitated during the slow cooling process after rolling to dissolve back into the matrix, this example subjected the rolled titanium alloy to a solution treatment followed by quenching. This treatment modulates the β-phase ratio and stacking fault energy, ultimately imparting the TRIP effect to the material, resulting in the preparation of an ultra-high-strength Ti6Al4V5Cu0.5Si0.25O titanium alloy exhibiting the TRIP effect.

[0077] The above description is merely an embodiment of the present invention and does not limit the scope of protection claimed by the present invention. Any equivalent structure or equivalent process transformation made using the contents of the present invention description and drawings, or directly or indirectly applied in other related technical fields, are also included in the scope of protection claimed by the present invention.

Claims

1. Ultra-high strength titanium alloy with TRIP effect, characterized by: The chemical composition and content of the titanium alloy meet the following requirements by weight percentage: Al：5.0~7.0； V: 3.0~5.0; Cu: 3.0~7.0; Si: 0.3~0.7; O: 0.15~0.35; the remainder is Ti.

2. The ultra-high strength titanium alloy with TRIP effect according to claim 1, characterized in that: The chemical composition and content of the titanium alloy meet the following requirements by weight percentage: Al：5.8~6.3； V: 3.7~4.2; Cu: 4.5~5.5; Si: 0.45~0.55; O: 0.22~0.27; the balance is Ti; The grade of ultra-high strength titanium alloy with TRIP effect is Ti6Al4V5Cu0.25Si0.25O.

3. The method for preparing the ultra-high-strength titanium alloy with TRIP effect according to claim 1, characterized in that: Step 1: using a vacuum consumable furnace to melt the ingot: the raw materials used for melting the ingot are required to be: Al: 5.0~7.0 according to the chemical composition and weight percentage; V: 3.0~5.0; Cu: 3.0~7.0; Si: 0.3~0.7; O: 0.15~0.35; the remainder is Ti; It is then forged into billets at temperatures above 1000°C; Step 2: keeping the forging blank at above 1000° C. for 1 to 3 hours and then quenching to obtain a fully martensitic structure; Step 3: by adding Cu, Si, and O elements, during the quenching process, Cu, Si, and O are supersaturated and dissolved in the β matrix, causing strong lattice distortion and refining the width of the martensite laths; the width of the martensite laths after quenching is refined to obtain an ultrafine lath martensite structure; Step 4, using the aforementioned ultrafine lath martensite as the initial structure, hot rolling the billet in the (α+β+Ti5Si3) three-phase region at 770-850°C to obtain a nano-ultrafine grain structure; Step 5: The rolled and deformed material is subjected to solution treatment at 800-840°C and then quenched, so that the Ti2Cu phase precipitated during the rolling cooling process dissolves back into the matrix, giving the material a TRIP effect.

4. The method for preparing an ultra-high-strength titanium alloy having a TRIP effect according to claim 3, characterized in that: During the rolling process in step 4, the deformation amount is required to be not less than 70%.

5. The method for preparing an ultra-high strength titanium alloy having TRIP effect according to claim 3 or 4, characterized in that: The rolling deformation temperature is 800~820℃, and the solution treatment temperature after rolling is 810~820℃.

6. The method for preparing an ultra-high strength titanium alloy having TRIP effect according to claim 5, characterized in that: During the rolling process of step 4, the rolling deformation is required to be no less than 80%.

7. The method for preparing an ultra-high strength titanium alloy having TRIP effect according to claim 5, characterized in that: In step 1, all materials, calculated by weight percentage, meet the following requirements: Al：5.8~6.3； V: 3.7~4.2; Cu: 4.5~5.5; Si: 0.45~0.55; O: 0.22~0.27; the remainder is Ti.

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