Method for improving mechanical property of TiAl alloy

Abstract

The invention discloses a method for improving the mechanical property of a TiAl alloy, and belongs to the technical field of TiAl alloys, the method comprises the steps that a TiAl alloy cast ingot is smelted, and the TiAl alloy cast ingot has a beta solidification and (alpha + gamma) two-phase region; carrying out thermal deformation treatment on the TiAl alloy cast ingot to obtain a thermal deformation alloy; carrying out heat treatment on the thermally deformed alloy in an (alpha + gamma) two-phase region, carrying out furnace cooling to a preset first temperature, and then carrying out air cooling to obtain a heat-treated alloy; and aging treatment is conducted on the heat treatment alloy at the preset second temperature, and the TiAl alloy is obtained. According to the method, thermal deformation and thermal treatment processes are combined, and the technical effect of improving the mechanical property of the TiAl alloy is achieved.

Description

Technical Field

[0001] This application relates to the technical field of TiAl alloys, and particularly to a method for improving the mechanical properties of TiAl alloys. Background Art

[0002] γ-TiAl alloys have unique advantages such as low density (3.8 - 4.2 g / cm 3 ), high specific strength and specific modulus, and good creep resistance and oxidation resistance, and are highly competitive lightweight high-temperature resistant materials for aero-engine systems. However, the poor room-temperature plasticity and difficult hot forming caused by their intermetallic compound properties have always been the key bottlenecks restricting their development and application.

[0003] Currently, the main TiAl alloy materials include cast alloys with high Al content (traditional γ-TiAl alloys) and β-solidified γ-TiAl alloys with low Al content. The fully lamellar structure in cast alloys is coarse and has low high-temperature strength. β-solidified γ-TiAl alloys will retain some β phase during the high-temperature process, and the high-temperature β phase will transform into β o phase orderly when it is retained to room temperature. The existence of this phase will damage the room-temperature plasticity of the alloy and affect the high-temperature performance of the alloy at the same time. Therefore, the existing TiAl alloy materials still have the problem of poor mechanical properties in high-temperature environments. Summary of the Invention

[0004] The main purpose of this application is to provide a method for improving the mechanical properties of TiAl alloys, aiming to solve the technical problems of poor strength and plasticity of TiAl alloys in high-temperature environments.

[0005] To achieve the above purpose, this application provides a method for improving the mechanical properties of TiAl alloys, including:

[0006] Melting a TiAl alloy ingot, where the TiAl alloy ingot has β-solidification and (α + γ) two-phase regions;

[0007] Performing hot deformation treatment on the TiAl alloy ingot to obtain a hot-deformed alloy;

[0008] Performing heat treatment on the hot-deformed alloy in the (α + γ) two-phase region, cooling it in the furnace to a preset first temperature and then air-cooling to obtain a heat-treated alloy;

[0009] Performing aging treatment on the heat-treated alloy at a preset second temperature to obtain a TiAl alloy.

[0010] In some embodiments of this application, the deformation amount of the hot deformation treatment is greater than 30%.

[0011] In some embodiments of the present application, the TiAl alloy ingot is held at 1300°C - 1380°C for 0.5 h - 1.0 h and then subjected to hot deformation treatment; the final deformation temperature of the hot deformation treatment is 1100°C - 1150°C.

[0012] In some embodiments of the present application, the steps of the hot deformation treatment are carried out in a non-isothermal atmospheric environment without a jacket.

[0013] In some embodiments of the present application, the preset first temperature is 1000°C - 1100°C.

[0014] In some embodiments of the present application, the cooling rate of furnace cooling is 0.15°C / s - 0.30°C / s.

[0015] In some embodiments of the present application, the preset second temperature is 800°C - 850°C; the aging treatment time is 3 h - 6 h.

[0016] In some embodiments of the present application, in atomic percentage, the main system of the TiAl alloy ingot is Ti-(40 - 44)Al-(1 - 3)Mn-(0 - 1.0)(Mo,W).

[0017] In some embodiments of the present application, in atomic percentage, the main system of the TiAl alloy ingot is Ti-(43 - 45)Al-(1 - 2)Mn-(2 - 3.5)Nb-(0 - 1)W.

[0018] In some embodiments of the present application, the process for melting the TiAl alloy ingot includes vacuum induction, vacuum induction + vacuum consumable, and plasma arc.

[0019] The method for improving the mechanical properties of TiAl alloy provided by the embodiments of the present application includes melting a TiAl alloy ingot, where the TiAl alloy ingot has β solidification and (α + γ) two-phase regions; performing hot deformation treatment on the TiAl alloy ingot to obtain a hot deformed alloy; performing heat treatment on the hot deformed alloy in the (α + γ) two-phase region, cooling it in the furnace to a preset first temperature and then air-cooling to obtain a heat-treated alloy; performing aging treatment on the heat-treated alloy at a preset second temperature to obtain a TiAl alloy. First, by utilizing the interlamellar stacking fault energy accumulated during the deformation of the alloy, a small number of γ-phase nucleation sites are created between the crystal clusters. The subsequent slow cooling process in the furnace provides relatively sufficient time and a high-temperature environment suitable for the growth of the nucleated γ phase. At the same time, since the equilibrium ratio of the γ phase in the (α + γ) two-phase region gradually increases as the temperature decreases, during the slow cooling process, the γ phase continuously nucleates and grows, and finally a certain number of spheroidized γ phases are formed between the crystal clusters. The formation of this spheroidized γ phase can not only pin the growth of α-phase grains, but also release the stacking fault energy brought by deformation by forming a new phase, balance the distribution of alloying elements, and enhance the tissue stability of the alloy during high-temperature service. This spherical γ phase at the grain boundary can deform by slip through slip systems, including ordinary dislocations 1 / 2<110] and superdislocations ((0 11) or 1 / 2<11 2), and can also generate twin 1 / 6<112](111) deformation through the slip of ordered twins, improving the room-temperature plasticity and high-temperature plasticity of the alloy. At the same time, due to the formation of small-sized spheroidized γ phases, the average crystal cluster size of the alloy decreases, resulting in a further increase in strength. This near fully lamellar structure with a small amount of spheroidized γ phases obtained by the synergistic regulation method of plastic deformation and heat treatment not only has higher strength, but also effectively improves its room-temperature and high-temperature plasticity. It has good hot workability and excellent tissue stability at the same time, and the mechanical properties are improved. BRIEF DESCRIPTION OF THE DRAWINGS

[0020] In order to more clearly illustrate the technical solutions in the embodiments of the present drawings or related technologies, the following will briefly introduce the drawings required for use in the description of the embodiments or related technologies. Obviously, the following drawings are only some embodiments of the present drawings. For those of ordinary skill in the art, without creative efforts, other drawings can also be obtained based on the structures shown in these drawings.

[0021] Figure 1 It is a schematic diagram of the heat treatment process for Example 1;

[0022] Figure 2 It is a microstructural diagram of the hot-rolled bar for Example 1;

[0023] Figure 3 It is a microstructural diagram of the undeformed as-cast alloy for Comparative Example 1;

[0024] Figure 4Microstructure diagram of the hot-rolled bar of Comparative Example 2;

[0025] Figure 5 Tensile strength test result diagrams of alloys in different states at room temperature, 750 °C, 800 °C, and 850 °C;

[0026] Figure 6 Elongation test result diagrams of alloys in different states at room temperature, 750 °C, 800 °C, and 850 °C;

[0027] Figure 7 Microstructure diagram of Example 1 after aging treatment at 800 °C for 1000 h;

[0028] Figure 8 Tensile test diagrams of the hot-rolled bar of Example 1 at room temperature, 750 °C, 800 °C, and 850 °C after aging for 1000 h.

[0029] The realization, functional features, and advantages of the object of the present application will be further described with reference to the embodiments and the accompanying drawings. Detailed implementation manners

[0030] It should be understood that the specific embodiments described herein are only used to explain the present application and are not used to limit the present application.

[0031] Hereinafter, the embodiments of the method for improving the mechanical properties of TiAl alloys disclosed in the present application will be specifically described in detail with appropriate reference to the accompanying drawings. However, there may be cases where unnecessary detailed descriptions are omitted. For example, there are cases where the detailed descriptions of well-known matters and the repeated descriptions of actually identical structures are omitted. This is to avoid the following description from becoming unnecessarily lengthy and to facilitate the understanding of those skilled in the art. In addition, the accompanying drawings and the following description are provided for those skilled in the art to fully understand the present application and are not intended to limit the subject matter recited in the claims.

[0032] The "range" disclosed in this application is defined in the form of a lower limit and an upper limit. A given range is defined by selecting a lower limit and an upper limit, and the selected lower limit and upper limit define the boundary of a particular range. The ranges defined in this way can include the end values or not include the end values, and can be combined arbitrarily, that is, any lower limit can be combined with any upper limit to form a range. For example, if ranges of 60-120 and 80-110 are listed for a specific parameter, ranges of 60-110 and 80-120 are also contemplated. In addition, if the minimum range values 1 and 2 are listed, and if the maximum range values 3, 4, and 5 are listed, then the following ranges are all contemplated: 1-3, 1-4, 1-5, 2-3, 2-4, and 2-5. In this application, unless otherwise specified, the numerical range "a-b" represents an abbreviated representation of any real number combination between a and b, where a and b are both real numbers. For example, the numerical range "0-5" means that all real numbers between "0-5" are fully listed herein, and "0-5" is only an abbreviated representation of these numerical combinations. In addition, when stating that a certain parameter is an integer ≥2, it is equivalent to disclosing that the parameter is, for example, the integers 2, 3, 4, 5, 6, 7, 8, 9, 10, 11, 12, etc.

[0033] If there is no special instruction, all the embodiments and optional embodiments of this application can be combined with each other to form a new technical solution.

[0034] To make the above objects, features, and advantages of this application more obvious and understandable, the technical solutions of this application will be further described below in conjunction with the drawings and embodiments. However, this application is not limited to the listed embodiments, and should also include any other well-known changes within the scope of the rights required by this application.

[0035] As a new type of lightweight high-temperature structural material, γ-TiAl alloy has unique advantages such as low density (3.9-4.2 g / cm 3 ), high specific strength, and high creep resistance. It is expected to replace some nickel-based superalloys in the temperature range of 600-900 °C to prepare certain aerospace structural components and rotating or reciprocating structural components of ground power systems, improve the thrust-to-weight ratio and fuel efficiency of the power system, and reduce greenhouse gas emissions. However, as an intermetallic compound, practical problems such as poor room-temperature plasticity, unstable microstructure, and extremely difficult hot forming limit the development and application of this alloy. Therefore, solving the above problems is an urgent need in the research field of TiAl alloys.

[0036] Different from traditional γ-TiAl alloys, β-solidifying γ-TiAl alloys can be obtained by reducing the Al content of the alloy (<45 at.%) and simultaneously adding sufficient amounts of β-stabilizing elements such as W, Mo, Mn, Nb, etc. Due to passing through the β single-phase region at high temperature, the β phase (A2, Im-3m) has a sufficient number of independent <111>{110} slip systems, enabling such alloys to exhibit certain hot working properties at high temperature. However, due to the strong action of β-phase stabilizing elements, a certain amount of the β phase in the alloy at high temperature cannot be completely transformed into the α phase during rapid cooling, resulting in the retention of a certain amount of the ordered βo phase in the room-temperature microstructure. The presence of this phase will damage the room-temperature plasticity of the alloy and simultaneously affect the high-temperature properties of the alloy. On the other hand, it is generally believed that the fully lamellar microstructure has good fracture toughness and creep resistance, and the remaining βo phase will also affect the final microstructure of the alloy and prevent the attainment of the fully lamellar microstructure.

[0037] To ensure that γ-TiAl alloys have both good hot workability and mechanical properties, a design concept for a new type of β-solidifying γ-TiAl alloy with both β-solidifying characteristics and the α single-phase region has been proposed. The ideal solidification route corresponding to this type of alloy is: L→L+β→β→β+α→α→α+γ→α2+(βo). There is a high-temperature β phase in this type of alloy system, and the sufficient number of slip systems in the high-temperature β phase ensures that the alloy has good hot deformation ability. In addition, the deformed alloy is heat-treated in the α single-phase region to eliminate the remaining βo phase and obtain a fully lamellar microstructure. In actual heat treatment, similar to traditional γ-TiAl alloys, due to the lack of pinning phases, the α-phase grains will significantly coarsen during heat treatment in the α single-phase region, resulting in large-sized fully lamellar grain clusters (200 - 1000 μm), leading to low strength and plasticity of the alloy and failing to achieve the designed intended purpose. To avoid the above situation, the heat treatment temperature of this new type of β-solidifying γ-TiAl alloy is selected in the (α+γ) two-phase region. During heat treatment, the undissolved γ phase can play a pinning role, not only avoiding the coarsening of α-phase grains but also ensuring that the final microstructure of the alloy is basically fully lamellar. However, since the cooling method adopted in the traditional heat treatment system is often air cooling with a relatively fast cooling rate, the γ phase precipitated in the (α+γ) two-phase region of the alloy and the elemental distribution are in an unbalanced state, resulting in the precipitation of fine γ-phase and α-phase mixed particle structures between the lamellar grain clusters during subsequent aging. Moreover, the volume fraction of this mixed structure will increase in the form of consuming the lamellae with the prolongation of the aging time, ultimately accelerating the degradation of the lamellar microstructure and damaging the service reliability of the alloy. In addition, although the fully lamellar microstructure obtained by the above heat treatment method has relatively high strength, its room-temperature plasticity is still poor. Therefore, to further improve the comprehensive properties of this new type of β-solidifying γ-TiAl alloy, methods that can simultaneously improve the strength and plasticity of the alloy and the microstructure stability need to be explored.

[0038] Based on this, the embodiments of the present application provide a method for improving the mechanical properties of TiAl alloys. In this embodiment, the method for improving the mechanical properties of TiAl alloys includes the following steps S10 to S40:

[0039] Step S10, melting a TiAl alloy ingot, the TiAl alloy ingot having β solidification and an (α + γ) two-phase region;

[0040] In the embodiments of the present application, by adjusting the alloy composition, the proportion of each element in the raw materials can be controlled, and a TiAl alloy ingot with β solidification and an (α + γ) two-phase region can be obtained by melting. Its solidification route is L → L + β → β → β + α → α → α2 + γ.

[0041] In some embodiments of the present application, the process for melting the TiAl alloy ingot includes vacuum induction, vacuum induction + vacuum consumable, and plasma arc. Vacuum induction refers to a process in which metal is melted by eddy currents generated by electromagnetic induction in a vacuum environment. Vacuum consumable means that in a vacuum, the material to be melted serves as one electrode, and the water-cooled copper crucible serves as the other electrode. An arc is drawn between the two electrodes, and the material to be melted is melted by the high temperature of the arc and drips into the crucible, gradually melting and gradually condensing into an ingot. Vacuum consumable can be combined with vacuum induction, that is, a comprehensive process of vacuum induction and vacuum consumable is used to melt the raw materials to obtain a TiAl alloy ingot. Plasma arc is a process in which a gas is used as a medium and a plasma arc or plasma beam is used as a heat source for melting, and is suitable for the production of superalloys. The above melting methods can be selected according to the actual situation.

[0042] In some embodiments of the present application, in atomic percentage, the main system of the TiAl alloy ingot is Ti-(40 - 44)Al-(1 - 3)Mn-(0 - 1.0)(Mo, W). For example, the main system of the TiAl alloy ingot can be Ti-42Al-1.5Mn-0.4Mo, Ti-42Al-1.5Mn-0.4W, or Ti-42Al-1.5Mn-0.2Mo-0.2W. Adding Mo or W to the Ti-Al-Mn series alloy endows the TiAl alloy with good tissue stability and oxidation resistance, and can be hot deformed under the conditions of no cladding and non-isothermal.

[0043] In some embodiments of the present application, in atomic percentage, the main system of the TiAl alloy ingot is Ti-(43-45)Al-(1-2)Mn-(2-3.5)Nb-(0-1)W. For example, the main system of the TiAl alloy ingot can be Ti-43Al-1.5Mn-3Nb or Ti-43Al-1.5Mn-3Nb-0.2W. Adding a small amount of Nb and W to the TiAl alloy can also improve the oxidation resistance of the TiAl alloy. At the same time, the Nb element helps to improve the microstructure of the TiAl alloy, enhancing the mechanical properties and processing performance of the TiAl alloy.

[0044] Step S20: Perform hot deformation treatment on the TiAl alloy ingot to obtain a hot deformed alloy;

[0045] The hot deformation treatment refers to a processing technology that deforms the TiAl alloy ingot at a certain temperature. During the hot deformation treatment, the grains in the TiAl alloy ingot are refined, and the TiAl alloy can accumulate the stacking fault energy between the lamellae, creating γ-phase nucleation points between the crystal clusters, providing a basis for the subsequent transformation of the γ-phase during heat treatment. The hot deformation treatment methods can include hot forging and rolling deformation.

[0046] In some embodiments of the present application, the deformation amount of the hot deformation treatment is greater than 30%. For example, the deformation amount of the hot deformation treatment can be 31%, 35%, 38%, 40%, 50%, 60%, 70%, 80%, etc. If the deformation amount is too small, the degree of grain refinement is small, which is less helpful for improving the alloy strength. A sufficient deformation amount can obtain a higher degree of grain refinement and accumulate a certain amount of stacking fault energy between the layers.

[0047] In some embodiments of the present application, the TiAl alloy ingot is held at 1300°C - 1380°C for 0.5h - 1.0h and then subjected to hot deformation treatment. For example, the holding temperature of the TiAl alloy ingot can be 1300°C, 1320°C, 1340°C, 1360°C, 1380°C or within the range composed of any of the above values. The holding time of the TiAl alloy ingot can be 0.5h, 0.6h, 0.7h, 0.8h, 0.9h, 1.0h or within the range composed of any of the above values. The holding treatment can make the TiAl alloy ingot reach a more uniform temperature distribution, providing good thermodynamic conditions for the subsequent hot deformation treatment.

[0048] In some embodiments of the present application, the final deformation temperature of the hot deformation treatment is 1100°C - 1150°C. For example, the final deformation temperature of the hot deformation treatment can be 1100°C, 1110°C, 1120°C, 1130°C, 1140°C, 1150°C, or within the range composed of any of the above values. The final deformation temperature of the hot deformation treatment can be adjusted according to the composition and phase transformation characteristics of the TiAl alloy. Under the alloy main system composition design of the embodiments of the present application, by setting an appropriate final deformation temperature, good microstructures and mechanical properties can be obtained.

[0049] In some embodiments of the present application, the steps of the hot deformation treatment are carried out in a non-isothermal atmospheric environment without a jacket. Through the regulation of the alloy composition, the TiAl alloy of the embodiments of the present application has good oxidation resistance, and there is no need to adopt a process with harsh conditions for jacket processing, and it can be directly processed in a non-isothermal atmospheric environment without a jacket.

[0050] Step S30: Heat-treat the hot-deformed alloy in the (α + γ) two-phase region, cool it in the furnace to a preset first temperature and then air-cool it to obtain a heat-treated alloy;

[0051] The heat treatment of the hot-deformed alloy in the (α + γ) two-phase region is carried out for heat preservation, and the heat treatment time can be set to 0.5 h - 1.0 h. The temperature of the (α + γ) two-phase region can be determined according to the phase transformation route of the alloy and is close to the γ-phase dissolution temperature. Furnace cooling refers to the process of cooling with the furnace. After the heat treatment, the cooling method of furnace cooling is adopted to slow down the cooling rate of the alloy. The slow cooling process can provide a relatively sufficient time and high-temperature environment suitable for the growth of the already nucleated γ-phase. At the same time, since the equilibrium ratio of the γ-phase in the (α + γ) two-phase region gradually increases with the decrease of temperature, during the slow cooling process, the γ-phase continuously nucleates and grows, and finally a certain amount of spheroidized γ-phase is formed between the crystal clusters. This spheroidized γ-phase can enhance the tissue stability of the alloy during the high-temperature service stage. The cooling rate of furnace cooling is slow. After cooling for a certain time until the temperature reaches the preset first temperature, enough spheroidized γ-phase can be obtained, and then it is air-cooled to room temperature. The preset first temperature is lower than the eutectoid temperature. A near fully lamellar structure with a certain amount of spheroidized γ-phase between the crystal clusters can be obtained, and the average lamellar cluster size can be controlled within 20 μm - 50 μm.

[0052] In some embodiments of the present application, the preset first temperature is 1000°C - 1100°C. For example, the preset first temperature can be 1000°C, 1020°C, 1040°C, 1060°C, 1080°C, 1100°C, or within the range composed of any of the above values. If the preset first temperature is too high, the cooling time is short, the number of spheroidized γ-phase is small, the pinning effect on the α-phase grains is weak, and the improvement of tissue stability is insufficient.

[0053] In some embodiments of the present application, the cooling rate of furnace cooling is 0.15 °C / s - 0.30 °C / s. For example, the cooling rate of furnace cooling can be 0.15 °C / s, 0.20 °C / s, 0.25 °C / s, 0.30 °C / s, or within the range composed of any of the above values. The cooling rate of furnace cooling cannot be too fast. If the cooling rate is too fast, the γ phase precipitated by the alloy in the (α + γ) two-phase region and the element distribution are in an unbalanced state, resulting in the precipitation of fine γ-phase and α-phase mixed particle structures between the lamellar crystal clusters during the subsequent aging process, and the volume fraction of this mixed structure will increase in the form of consuming the lamellae with the extension of the aging time, ultimately accelerating the degradation of the lamellar structure and damaging the service reliability of the alloy.

[0054] Step S40: Perform aging treatment on the heat-treated alloy at a preset second temperature to obtain a TiAl alloy.

[0055] Aging treatment refers to a heat treatment process in which a metal or alloy workpiece undergoes a certain degree of cold working deformation and then is placed at a higher temperature or room temperature to maintain its shape, size, and its properties change with time.

[0056] In some embodiments of the present application, the preset second temperature is 800 °C - 850 °C; the time of aging treatment is 3 h - 6 h. For example, the preset second temperature of aging treatment can be 800 °C, 810 °C, 820 °C, 830 °C, 840 °C, 850 °C, or within the range composed of any of the above values. The time of aging treatment can be 3 h, 3.5 h, 4 h, 4.5 h, 5 h, 5.5 h, 6 h.

[0057] In this embodiment, a TiAl alloy ingot is melted, and the TiAl alloy ingot has a β solidification and (α+γ) two-phase region; the TiAl alloy ingot is subjected to heat deformation treatment to obtain a heat-deformed alloy; the heat-deformed alloy is heat-treated in the (α+γ) two-phase region, furnace-cooled to a preset first temperature, and then air-cooled to obtain a heat-treated alloy; the heat-treated alloy is subjected to aging treatment at a preset second temperature to obtain a TiAl alloy. First, the interlamellar stacking fault energy accumulated by the alloy during the deformation process is used to create a small amount of γ phase nucleation points between the crystal clusters. The subsequent slow cooling process with the furnace provides the nucleated γ phase with relatively sufficient time and high temperature environment suitable for growth. At the same time, since the equilibrium ratio of the γ phase in the (α+γ) two-phase region gradually increases with decreasing temperature, the γ phase continuously nucleates and grows during the slow cooling process, and finally forms a certain number of spheroidized γ phases between the crystal clusters. The formation of the spheroidized γ phase can not only pin the growth of α phase grains, but also release the stacking fault energy brought by deformation by forming new phases, balance the distribution of alloy elements, and enhance the structural stability of the alloy in the high-temperature service stage. This spheroidal γ phase at the grain boundary can not only deform through slip system, including ordinary dislocation 1 / 2<110] and super dislocation ((0 11) or 1 / 2<11 2), but also produce twin 1 / 6<112](111) deformation through the slip of ordered twins, which improves the room temperature plasticity and high temperature plasticity of the alloy. At the same time, due to the formation of small-sized spheroidized γ phase, the average crystal size of the alloy is reduced, which further improves the strength. The nearly full lamellar structure with a small amount of spheroidized γ phase obtained by the coordinated regulation of plastic deformation and heat treatment not only has higher strength, but also effectively improves its room temperature and high temperature plasticity. It has good hot workability and excellent structural stability, and its mechanical properties are improved.

[0058] The present application scheme is described in detail below with reference to specific embodiments.

[0059] Example 1

[0060] The main component of the alloy is Ti-43Al-1.5Mn-3Nb (at.%). The phase transformation route of the alloy is calculated using PANDATTM (2023) thermodynamic calculation software, and the solidification route of the alloy is: L→L+β→β→β+α→α→α2+γ, which satisfies the β solidification characteristics and has a (α+γ) two-phase region. Through the analysis of the alloy phase diagram, the γ phase dissolution temperature of the alloy is approximately equal to 1250°C.

[0061] The preparation method of alloy ingots includes: the main raw materials for alloy preparation are sponge titanium, industrial pure aluminum, purified manganese, and aluminum-niobium master alloy. The ingredients are prepared according to the main components of the alloy, 20 kg of alloy material is melted in a vacuum induction melting furnace, and cast into 4 ingots of size Alloy ingots. The alloy ingots are directly rolled into bars with a diameter of 12 mm by a Y-type rolling mill in multiple passes at one time, with a deformation amount > 70%, and the initial heating temperature for rolling is 1380 °C.

[0062] The rolled bars are subjected to a high-temperature treatment at 1220 °C (the temperature of the (α + γ) two-phase region) for 0.5 h. After the treatment is completed, they are slowly cooled in the furnace at a cooling rate of 0.19 °C / s to 1000 °C, taken out and air-cooled (AC) to room temperature. Finally, an aging treatment is carried out at 850 °C for 3 h, and the cooling method after aging is furnace cooling (FC). The heat treatment flow chart is as Figure 1 shown.

[0063] Comparative Example 1

[0064] The main components of the alloy are the same as those in Example 1, and the preparation method of the alloy ingots is the same as that in Example 1. However, no deformation treatment is carried out. The undeformed as-cast alloy is subjected to a high-temperature treatment at 1220 °C (the (α + γ) two-phase region) for 0.5 h. After the treatment is completed, it is cooled in the furnace at a cooling rate of 0.19 °C / s to 1000 °C, taken out and air-cooled (AC) to room temperature. Finally, an aging treatment is carried out at 850 °C for 3 h, and the cooling method after aging is furnace cooling (FC).

[0065] Comparative Example 2

[0066] The main components of the alloy are the same as those in Example 1, and the preparation method of the alloy ingots is the same as that in Example 1. The same rolling deformation treatment as in Example 1 is also carried out. The rolled bars are subjected to a high-temperature treatment at 1220 °C (the (α + γ) two-phase region) for 0.5 h. After the treatment is completed, they are directly taken out and air-cooled (AC) to room temperature. Finally, an aging treatment is carried out at 850 °C for 3 h, and the cooling method after aging is furnace cooling (FC).

[0067] Alloy microstructure analysis: The polished microstructures of the series of treated specimens in Example 1, Comparative Example 1, and Comparative Example 2 are taken, and an electron probe with the model JXA-8530F is used to observe each polished microstructure under the condition of backscattered electrons.

[0068] Figure 2 is the microstructural diagram of the hot-rolled bar in Example 1. Figure 3 is the microstructural diagram of the undeformed as-cast alloy in Comparative Example 1. Figure 4 is the microstructural diagram of the hot-rolled bar in Comparative Example 2. Combining Figures 2 to 4 It can be seen that the microstructure obtained after slow-cooling heat treatment of the deformed rolled bar is small-sized (α2 / γ) lamellar crystal clusters, and spherical γ is accompanied between the crystal clusters. gPhase; while the as-cast alloy without deformation only obtains a fully lamellar structure with larger grain cluster sizes after the same slow cooling heat treatment. It should be noted that there will be no (α2 + γ) mixed heterophase as shown in Comparative Example 2 between the lamellar grain clusters of the structure treated by the slow cooling heat treatment system, indicating that introducing the slow cooling heat treatment system is beneficial to improving the tissue stability.

[0069] Tensile property analysis of the alloy: Standard tensile specimens were machined from samples taken from the as-cast structure after slow cooling heat treatment (Comparative Example 1), hot-rolled bars, hot-rolled bars air-cooled treatment (Comparative Example 2), and hot-rolled bars slow cooling heat treatment (Example 1), and tensile tests were carried out on a tensile testing machine at room temperature, 750 °C, 800 °C, and 850 °C to evaluate their comprehensive mechanical properties (the room temperature and high-temperature tensile properties were respectively carried out in accordance with the standards of GB / T228.1-2010 and GB / T228.2-2015).

[0070] Figure 5 It is a graph of the tensile strength test results of alloys in different states at room temperature, 750 °C, 800 °C, and 850 °C. Figure 6 It is a graph of the elongation test results of alloys in different states at room temperature, 750 °C, 800 °C, and 850 °C. Combining Figure 5 and Figure 6 It can be seen that the near-fully lamellar structure obtained by the slow cooling heat treatment system for the alloy with a certain amount of deformation in Example 1 has the best strength at each tensile temperature. Alloys that do not adopt the slow cooling heat treatment system or adopt the slow cooling heat treatment system but have no initial deformation have almost no plasticity at room temperature and poor high-temperature plasticity; however, the alloy treated by the synergistic treatment of plastic deformation and slow cooling heat treatment system not only improves the room temperature plasticity, but also greatly improves the high-temperature plasticity.

[0071] Analysis of the tissue stability of the alloy: After slow cooling heat treatment (Example 1) of the rolled bar, aging treatment was carried out at 800 °C for 1000 h. Figure 7 It is the microstructure diagram of Example 1 after aging treatment at 800 °C for 1000 h. Observing Figure 7 the tissue after aging treatment, it is found that there is no obvious difference in the tissue of the alloy obtained by the slow cooling heat treatment system before and after aging, indicating that the tissue has high stability. Tensile tests were carried out on the rolled bar after 1000 h of aging at room temperature, 750 °C, 800 °C, and 850 °C, and the results are shown in Figure 8 . It can be found that long-term high-temperature aging treatment only slightly reduces its strength, while the plasticity is further improved, which further verifies that the proposed synergistic treatment method of plastic deformation and slow cooling heat treatment can not only improve the room temperature and high-temperature mechanical properties of the alloy, but also optimize the tissue and performance stability of the alloy during the high-temperature service stage.

[0072] It should be noted that this application is not limited to the above-described embodiments. The above-described embodiments are merely examples, and embodiments having the same composition and achieving the same effects as the technical idea within the scope of the technical solution of this application are all included in the technical scope of this application. In addition, within the scope of not departing from the gist of this application, various modifications that can be conceived by those skilled in the art to the embodiments, and other ways constructed by combining some of the constituent elements in the embodiments are also included in the scope of this application.

Claims

1. A method for improving the mechanical properties of TiAl alloy, characterized in that, Including: Smelting a TiAl alloy ingot, the TiAl alloy ingot having β solidification and an (α+γ) two-phase region; Performing hot deformation treatment on the TiAl alloy ingot to obtain a hot deformed alloy; Performing heat treatment on the hot deformed alloy in the (α+γ) two-phase region, furnace cooling to a preset first temperature and then air cooling to obtain a heat treated alloy; Performing aging treatment on the heat treated alloy at a preset second temperature to obtain a TiAl alloy.

2. The method for improving the mechanical properties of a TiAl alloy according to claim 1, wherein The deformation amount of the hot deformation treatment is greater than 30%.

3. The method for improving the mechanical properties of the TiAl alloy according to claim 2, wherein The TiAl alloy ingot is held at 1300°C - 1380°C for 0.5 h - 1.0 h and then subjected to hot deformation treatment; the final stage deformation temperature of the hot deformation treatment is 1100°C - 1150°C.

4. The method for improving the mechanical properties of the TiAl alloy according to claim 3, characterized in that, The steps of the hot deformation treatment are carried out in a non-sheathed and non-isothermal atmospheric environment.

5. The method for improving the mechanical properties of TiAl alloy according to claim 1, characterized in that, The preset first temperature is 1000°C - 1100°C.

6. The method for improving the mechanical properties of the TiAl alloy according to claim 5, characterized in that, The cooling rate of the furnace cooling is 0.15°C / s - 0.30°C / s.

7. The method for improving the mechanical properties of the TiAl alloy according to claim 1, characterized in that, The preset second temperature is 800°C - 850°C; the time of the aging treatment is 3 h - 6 h.

8. The method for improving the mechanical properties of a TiAl alloy according to claim 1, characterized in that, In atomic percentage, the main system of the TiAl alloy ingot is Ti-(40 - 44)Al-(1 - 3)Mn-(0 - 1.0)(Mo,W).

9. The method for improving the mechanical properties of a TiAl alloy according to claim 1, characterized in that, In atomic percentage, the main system of the TiAl alloy ingot is Ti-(43 - 45)Al-(1 - 2)Mn-(2 - 3.5)Nb-(0 - 1)W.

10. The method for improving the mechanical properties of the TiAl alloy according to any one of claims 1 to 9, characterized in that, The process for smelting the TiAl alloy ingot includes vacuum induction, vacuum induction + vacuum consumable, and plasma arc.

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