Preparation method for nickel-based superalloy

Abstract

The present application belongs to the technical field of superalloys. Disclosed is a preparation method for a nickel-based superalloy. The present application solves the problem of a nickel-based superalloy in the prior art making it difficult to satisfy the comprehensive requirements of a component for both a long alloy endurance life and a low crack propagation rate. The preparation method comprises: step 1: performing smelting, so as to obtain a cast ingot; step 2: subjecting the cast ingot to a homogenizing heat treatment, wherein the homogenizing heat treatment comprises two sections of heat preservation; and after the second-section heat preservation, performing furnace cooling to 1000+ / -10°C or below, and then performing discharging and air cooling; step 3: preparing a bar blank and a forging piece by means of forging; step 4: performing a solution treatment; and step 5: performing a stabilization and aging treatment, so as to obtain a nickel-based superalloy. The nickel-based superalloy prepared by means of the method of the present application has a high strength, good low-cycle fatigue performance, a low crack propagation rate and excellent comprehensive performance.

Description

A method for preparing nickel-based high-temperature alloy Technical Field

[0001] The present application relates to the technical field of high-temperature alloys, and in particular to a method for preparing a nickel-based high-temperature alloy. Background Art

[0002] Currently, aerospace annular components and disc forgings, such as gas turbine discs and blades, require materials that take into account both high-temperature durability and creep fatigue. The materials must ensure a certain low level of crack growth rate and have a high damage tolerance to ensure the service reliability of the materials. However, the commonly used nickel-based alloys currently have a fast crack growth rate and cannot meet the comprehensive requirements of some aerospace components for high alloy durability and low crack growth rate. Summary of the Invention

[0003] In view of the above situation, the present application aims to provide a method for preparing a nickel-based high-temperature alloy, which is used to solve the problem that existing nickel-based high-temperature alloys are difficult to simultaneously meet the comprehensive requirements of components for high alloy durability and low crack growth rate.

[0004] The purpose of this application is mainly achieved through the following technical solutions:

[0005] The present application provides a method for preparing a nickel-based high-temperature alloy, comprising:

[0006] Step 1: Smelting to obtain ingot;

[0007] Step 2: subject the ingot to homogenization heat treatment; the homogenization heat treatment includes two stages of heat preservation; after the second stage of heat preservation, the furnace is cooled to below 1000±10℃ and then taken out of the furnace for air cooling;

[0008] Step 3: Forging to prepare billets and forgings;

[0009] Step 4: solution treatment;

[0010] Step 5: Stabilization and aging treatment to obtain nickel-based high-temperature alloy.

[0011] Furthermore, in step 2, the process steps of homogenizing heat treatment include:

[0012] S201, slowly raise the temperature from ≤400℃ to 1165~1170℃, heating time 10~15h, keeping warm for 35~37h;

[0013] S202. Continue to heat up to 1200-1210℃, keep warm for 38-42h, then cool the furnace to below 1000±10℃ and take it out of the furnace for air cooling.

[0014] Furthermore, in step 4, the solution treatment process is: heating to 995-1020° C. and keeping the temperature for 3-5 hours, then taking the product out of the furnace and cooling it.

[0015] Furthermore, in step 4, the heating rate is 3 to 6°C / min.

[0016] Furthermore, in step 4, for disk-shaft parts with a wall thickness of 20 to 200 mm, cooling oil is used for cooling after leaving the furnace; for ring parts with a wall thickness of 10 to 100 mm, air cooling is used after leaving the furnace.

[0017] Furthermore, in step 5, the stabilization and aging treatment process is as follows: heating to 820-850°C and keeping it for 5-7 hours, air cooling to 500±10°C, continuing to heat to 740-770°C and keeping it for 8-10 hours, and then air cooling to room temperature.

[0018] Furthermore, in step 5, the heating rate is controlled to be 2-3.5°C / min.

[0019] Furthermore, the components of the nickel-based high-temperature alloy include, by mass percentage, C: 0.02%~0.06%, Cr: 18.5%~20.0%, Co: 13.0%~14.0%, Mo: 4.0%~4.90%, Al: 1.3%~1.6%, Ti: 2.80%~3.25%, B: 0.015%~0.03%, Zr: 0.05%~0.15%, Ti / Al: 2.25~2.38, (Al+Ti): 4.35%~4.58%, Fe: 0.05%~2.0%, Mg≤0.005%, P: 0.004%~0.009%, O: ≤20PPm, N: ≤20PPm, S≤10PPm, nickel: balance.

[0020] Furthermore, the microstructure of the nickel-based high-temperature alloy obtained in step 5 mainly includes equiaxed austenite grains, uniformly dispersed borides and carbides, and γ' phase dispersed inside the grains.

[0021] Furthermore, in the microstructure of nickel-based high-temperature alloys, carbides mainly include M 23 C6 and MC.

[0022] Compared with the prior art, this application can achieve at least one of the following beneficial effects:

[0023] a) In the preparation method of the nickel-based high-temperature alloy of the present application, by precisely controlling the process parameters of each step, for example, the homogenization heat treatment includes two stages of heat preservation; after the second stage of heat preservation, the furnace is cooled to below 1000±10°C and then air-cooled after being taken out of the furnace to reduce the cracking tendency; combined with precise control of other steps and the parameters of each step, deformation is avoided, the alloy grains are ensured to be uniform and fine, and the morphology and distribution of carbides in the alloy and the content and size of other second phases are ensured to be optimally matched, thereby achieving the goals of improving the alloy strength, reducing the cracking tendency, and reducing the crack growth rate, thereby ensuring the comprehensive performance of the nickel-based high-temperature alloy.

[0024] b) In the preparation method of the nickel-based high-temperature alloy of the present application, the content of individual elements such as B, Zr, Fe, C, Cr, Co, Al, and Ti in the alloy is precisely controlled to enhance the solid solution strengthening effect of the alloy and the grain boundary strength of the alloy; and by synergistically controlling the values ​​of B, Zr, and C, the optimal matching of the morphology and distribution of carbides and the content and size of other second phases in the alloy is ensured. Without reducing the strength of the alloy, the risk sources of cracks in the alloy are reduced, the carbides are uniformly distributed at the grain boundaries of the alloy, the grain boundary strength is enhanced, and the speed of dislocation slip is slowed, thereby achieving the purpose of improving the strength of the alloy, reducing the cracking tendency, and reducing the crack propagation rate, thereby ensuring the comprehensive performance of the alloy.

[0025] c) The properties of the nickel-based high-temperature alloy prepared by the preparation method of the present application are as follows: Room temperature properties: tensile strength σ b ≥1300MPa (e.g. 1330-1360MPa); yield strength σ 0.2 ≥1050MPa (e.g., 1052-1080MPa); elongation after fracture δ5 ≥27% (e.g., 27%-28%); cross-sectional shrinkage ψ ≥35% (e.g., 35%-40%); 730℃ / 550MPa endurance performance: endurance time τ ≥70h (e.g., 72-88h); elongation after fracture δ5 ≥17.5% (e.g., 17.7%-20%); 815℃ / 295MPa endurance performance: endurance time τ ≥83h (e.g., 83-95h); elongation after fracture δ5 ≥26% (e.g., 26%-30%); low-cycle fatigue performance: 500℃ / strain control 0-0.7% / 0.33Hz, >5×10 4 Week (e.g., 57532 to 63645 weeks).

[0026] Other features and advantages of the present application will be described in the subsequent description, and some will become apparent from the description or be understood by practicing the present application. The objectives and other advantages of the present application can be realized and obtained by the contents particularly pointed out in the written description, claims, and drawings. BRIEF DESCRIPTION OF THE DRAWINGS

[0027] The accompanying drawings are only for the purpose of illustrating particular embodiments and are not to be considered as limiting the present application. Like reference symbols denote like components throughout the drawings.

[0028] FIG1 is a graph showing crack growth rates of Example 1 and Comparative Example 4;

[0029] FIG2 is one of the microstructure diagrams of Example 1 of the present application;

[0030] FIG3 is a second microstructure diagram of Example 1 of the present application;

[0031] FIG4 is a microstructure diagram of Comparative Example 4;

[0032] FIG5 is a grain map of Comparative Example 4. DETAILED DESCRIPTION

[0033] The preferred embodiments of the present application are described in detail below in conjunction with the accompanying drawings, wherein the accompanying drawings constitute a part of the present application and are used together with the embodiments of the present application to illustrate the principles of the present application.

[0034] The present application provides a method for preparing a nickel-based high-temperature alloy, comprising:

[0035] Step 1: Smelting to obtain ingot;

[0036] Step 2: subject the ingot to homogenization heat treatment; the homogenization heat treatment includes two stages of heat preservation; after the second stage of heat preservation, the furnace is cooled to below 1000±10℃ and then taken out of the furnace for air cooling;

[0037] Step 3: Forging to prepare billets and forgings;

[0038] Step 4: solution treatment;

[0039] Step 5: Stabilization and aging treatment to obtain nickel-based high-temperature alloy.

[0040] Specifically, the components of the above-mentioned nickel-based high-temperature alloy include, by mass percentage, C: 0.02%~0.06%, Cr: 18.5%~20.0%, Co: 13.0%~14.0%, Mo: 4.0%~4.90%, Al: 1.3%~1.6%, Ti: 2.80%~3.25%, B: 0.015%~0.03%, Zr: 0.05%~0.15%, Ti / Al: 2.25~2.38, (Al+Ti): 4.35%~4.58%, Fe: 0.05%~2.0%, Mg≤0.005%, P: 0.004%~0.009%, O: ≤20PPm, N: ≤20PPm, S≤10PPm, nickel: balance.

[0041] Specifically, in the above step 1, VIM+ESR+VAR triple smelting or VIM+VAR double smelting is adopted to form ingots.

[0042] Specifically, in step 2 above, the purpose of the homogenization heat treatment is to eliminate the low-melting-point phase and eutectic phase in the ingot and reduce element segregation. The homogenization heat treatment is set up with two stages of heat preservation. In the first stage of heat preservation, the temperature is lower and the main function is to eliminate the low-melting-point phase and eutectic phase in the alloy; the second stage of heat preservation can promote the uniform diffusion of segregated elements. Specifically, the process steps of the homogenization heat treatment include:

[0043] S201, heat to 1165-1170℃, keep warm for 35-37h;

[0044] S202. Continue to heat up to 1200-1210℃, keep warm for 38-42h, then cool the furnace to below 1000±10℃ and take it out of the furnace for air cooling.

[0045] Specifically, in S201, the temperature is slowly increased from a furnace temperature of ≤400° C. to 1165-1170° C., and the heating time is 10-15 hours.

[0046] Specifically, in S201, considering that excessively high temperatures and prolonged holding times will cause the low-melting-point phase and eutectic phase formed by the B element in the as-cast structure to remelt, this will, on the one hand, form pores after homogenization, hindering subsequent alloy cogging and even causing the ingot to crack and become scrapped. Furthermore, the B element will not be able to fully diffuse into the matrix to prepare for the subsequent formation of grain boundary strengthening phases, ultimately failing to strengthen the grain boundaries and improve the alloy's creep resistance. Therefore, the temperature is controlled to 1165-1170°C and held for 35-37 hours.

[0047] Specifically, in S202, the higher the temperature and the longer the holding time, the more complete the remelting of segregated elements. However, once the temperature and time reach a matching equilibrium point, the remelting of segregated elements will remain stable. Therefore, excessively high temperatures and long holding times have little positive effect on the remelting of segregated elements, and instead lead to coarse grains, energy waste, and reduced production efficiency. On the other hand, low temperatures and short holding times cannot ensure the majority of segregated elements are remelted, and dendritic segregation will reduce hot working plasticity during forging. Therefore, the temperature is controlled to 1200-1210°C and held for 38-42 hours.

[0048] Specifically, in S202, considering that more grain boundary strengthening elements and carbide-forming elements are added to the alloy of the present application, while the strength of the alloy is improved, the second phase is also increased, which increases the risk of cracking of the alloy during the ingot casting process. Therefore, the furnace is controlled to cool to 1000±10℃ and then air-cooled after being taken out of the furnace. The furnace cooling treatment can make the curved grain boundaries account for a larger proportion and the high-temperature plasticity better. Therefore, the homogenization treatment is controlled to increase the furnace cooling, improve the plastic toughness of the ingot during the casting process, and avoid cracking.

[0049] Specifically, in the above step 3, the forging method includes but is not limited to free forging of the rapid forging method or the combination of upsetting and drawing + radial forging method of blanking and subsequent hot rolling and cold drawing, wherein the rapid forging method includes but is not limited to single drawing and the combination of upsetting and drawing. The bar specifications of blanking forging or subsequent rolling and cold drawing can reach scope.

[0050] Specifically, in step 3 above, for stretching alone, the stretching deformation is 30-60% per fire, and the target size is generally achieved after three to four stretching cycles. For upsetting plus stretching, the upsetting deformation is 30-60% per fire, and the stretching deformation is 30-50% per fire. The bar production temperature range is 1060°C to 1200°C. The temperature is lowered by 40-80°C with each fire until the final fire temperature reaches below 1080°C. The upsetting and stretching cycles are performed in one cycle before the temperature is lowered.

[0051] Specifically, in the above step 3, the rod blank can be die-forged to produce a turbine disk with an integrated disk and shaft, or a ring-shaped part can be produced by ring rolling.

[0052] Specifically, in the above step 4, the solution treatment process is: heating to 995-1020° C. and keeping the temperature for 3-5 hours, then taking the product out of the furnace and cooling it.

[0053] Specifically, in step 4 above, considering that a rapid heating rate would prevent the alloy core from reaching the solution temperature, resulting in insufficient holding time, a slow heating rate of 3-6°C / min was used. Excessively high temperatures or long holding times would lead to grain growth, which would be detrimental to alloy properties. On the other hand, lower temperatures or short holding times would not allow for the formation of the appropriate γ' phase during the subsequent aging process. Therefore, the temperature was raised to 995-1020°C and held for 3-5 hours.

[0054] Specifically, in step 4 above, to prevent excessive stress and deformation caused by excessive cooling rates, different cooling methods are used for parts with different wall thicknesses. For example, for disc-shaft parts with a wall thickness of 20-200mm, cooling is performed by injecting cooling oil into the parts after leaving the furnace; for ring parts with a wall thickness of 10-100mm, air cooling is used after leaving the furnace.

[0055] Specifically, in the above step 5, the stabilization and aging treatment process is as follows: heating to 820-850°C and keeping it warm for 5-7 hours, air cooling to 500±10°C, continuing to heat to 740-770°C and keeping it warm for 8-10 hours, and then air cooling to room temperature.

[0056] Specifically, in step 5 above, consider that a rapid heating rate will prevent the alloy core from reaching the stabilization and aging temperature, resulting in insufficient holding time, insufficient precipitation of the γ' phase, and suboptimal quantity and size. Excessively high temperatures or prolonged holding times will result in larger γ' phases, deviating from the optimal size and quantity ratio, and affecting alloy strength. Therefore, the temperature is raised slowly, starting at a furnace temperature ≤ 400°C, with a controlled heating rate of 2-3.5°C / min.

[0057] Specifically, in step 5, the temperature is kept at 820-850℃ for 5-7 hours and then air-cooled to 500±10℃ instead of air-cooling to room temperature. This can improve efficiency and effectively ensure the size of carbides. For example, the size of MC carbides is 5-10μm; 23 The size of C6 carbide is about 0.3 to 1 μm.

[0058] Specifically, in the composition of the nickel-based high-temperature alloy, O+N+S≤50PPm, for example, O+N+S≤40PPm.

[0059] The following is a detailed description of the functions and dosages of the components of the nickel-based high-temperature alloy in this application:

[0060] C: As the C content increases, the carbides (MC, M 23 C6) will also increase, but M 23 The precipitation temperature of C6 does not change much. In this application, by controlling the C content, the beneficial phase M is precipitated at the grain boundary with the B element. 23 C6, so that it is distributed in a continuous or necklace shape, avoiding film-like precipitation. Continuous film-like precipitation is harmful to the alloy properties, while continuous or necklace-like precipitation is beneficial to the alloy properties. By appropriately reducing the C content and controlling the content of alloy impurity elements, the content of primary carbides and nitrides in the grains is effectively reduced, the fatigue cracking crack source is reduced, and the fatigue properties of the alloy are improved. Generally, when the C content is greater than 0.06%, the alloy's rupture life will be reduced. However, by increasing the content of B, Zr, and Mg, the grain boundary segregation of the B element makes the grain boundary film M 23 C6 is not easy to precipitate, which improves the grain boundary. Therefore, in this application, the C content is controlled to 0.02% to 0.06%.

[0061] Cr: Cr element can improve the oxidation resistance and corrosion resistance of the alloy. 23 The main forming elements of C6, its content is related to M 23 The content of C6 has little effect, but significantly affects M 23 The precipitation temperature of C6, the more Cr content, the higher the M 23 The higher the precipitation temperature of C6, when the Cr content exceeds 20%, M 23The precipitation temperature of C6 exceeds 960° C. Therefore, the Cr content is controlled within a range of 18.5% to 20.0% in this application.

[0062] Co: The Co element mainly plays a role in solid solution strengthening. The addition of Co to the γ matrix can reduce the stacking fault energy of the matrix. The reduced stacking fault energy increases the probability of stacking faults, making cross-slip of dislocations more difficult. In this way, deformation requires greater external force, which is manifested as an increase in strength. Moreover, the stacking fault energy is reduced, the creep rate is reduced, and the creep resistance is increased. In addition, the Co element can also reduce the solubility of the γ′-forming elements Ti and Al in the matrix, thereby increasing the amount of γ′ precipitation phase in the alloy and increasing the service temperature of the alloy. However, Co is a scarce resource in my country. Considering the cost factor of the alloy, this application controls the Co element content to 13.0% to 14.0%.

[0063] Mo: Mo enters the nickel-based alloy matrix, primarily providing solid solution strengthening. It also increases the matrix's lattice constant, reduces the γ-γ′ mismatch, and improves the alloy's stability and high-temperature creep properties. However, the addition of Mo lowers the alloy's initial and final melting temperatures. Therefore, the Mo content in this application is controlled within a range of 4.0% to 4.9%.

[0064] Al: Al is the primary element that forms the γ′ phase. Approximately 20% of the Al added to the alloy enters the γ solid solution, providing solid solution strengthening. The remaining 80% reacts with Ni to form Ni3Al, providing precipitation strengthening. To ensure the presence of the γ′ phase necessary to maintain high-temperature strength at 700°C, the Al content in this application is controlled between 1.3% and 1.6%.

[0065] Ti: When added to nickel-based alloys, approximately 10% of Ti enters the γ solid solution, providing a certain degree of solid solution strengthening, while approximately 90% enters the γ′ phase. Under certain Al content conditions, increasing Ti content increases the amount of γ′ phase and the alloy's high-temperature strength. To ensure the presence of the γ′ phase necessary to maintain high-temperature strength at 700°C, this application controls the Ti content to 2.80% to 3.25%, while also controlling the Ti / Al ratio to 2.25 to 2.38 and the (Al+Ti) ratio to 4.35% to 4.58%.

[0066] Increasing the Al or Ti content alone or simultaneously does not necessarily increase the secondary γ′ phase content at the corresponding solution temperature. However, a high (Al+Ti) content has a clear advantage because its primary γ′ phase has a higher dissolution temperature. In the temperature range where the primary γ′ phase in the low (Al+Ti) group is completely dissolved, while the primary γ′ phase in the high (Al+Ti) group is not completely dissolved, the secondary γ′ phase content of the high (Al+Ti) content is higher. Due to the pinning effect of the primary γ′ phase, reasonable temperature control can hinder the growth of grain size. At the same time, the two sizes of γ′ phase can play a joint strengthening role.

[0067] B: Boride (mainly including M3B2, M4B3, MB 12 ) The boride is evenly distributed in the grain through homogenization and cogging. By increasing the B content, the segregation of B atoms at the grain boundary is increased, the grain boundary bonding force is increased, and the activation energy of grain boundary diffusion is increased. The thermodynamic transformation is used to form a beneficial continuous or necklace-shaped M at the grain boundary. 23 C6 carbides can increase the surface energy of grain boundary cracks while hindering grain boundary sliding, thereby improving the grain boundary strength of the alloy and enhancing high-temperature creep performance. This significantly increases the creep limit of the material, especially its endurance strength. The higher the B content, the better the high-temperature creep performance. However, when the B content is high, more layered low-melting-point phases will form, which is detrimental to the improvement of alloy performance and instead leads to reduced performance. Therefore, the B content should not be too high. Taking all factors into consideration, the B content is controlled within the range of 0.015% to 0.03% in this application.

[0068] Fe: The lattice constant of Fe differs by 3% from that of Ni. Its lattice expansion induces long-range stress fields, hindering dislocation motion. Fe also reduces stacking fault energy, increases yield strength, and acts as a solid solution strengthener. The nickel-based superalloy in this application is primarily composed of grain boundary strengthening elements. While these elements are added, a portion of Fe is also incorporated to enhance matrix strength and achieve matching strength. Therefore, the Fe content in this application is controlled to 0.05% to 2.0%.

[0069] P: This application achieves beneficial P precipitation at grain boundaries by rationally controlling the P content, improving the alloy's high-temperature creep performance without compromising its plasticity. However, high P content can significantly increase notch sensitivity during high-temperature creep in some alloys. Due to the addition of appropriate amounts of B and Zr, the P content in this application is controlled to 0.004% to 0.009%.

[0070] Zr: As a grain boundary strengthening element, Zr can effectively reduce the solidus temperature of nickel-based alloys and thus expand the solidification range of the alloy. At the same time, it can promote the segregation of C and S elements at the grain boundaries together with Mg elements, purify the alloy, and the addition of an appropriate amount of Zr can improve the creep resistance of the alloy. However, too high a Zr content will form more eutectic phases such as Ni5Zr, which will have an adverse effect on the plasticity of the alloy. This application increases the content of Zr while increasing the content of B elements. The strength of the material is improved by the segregation of Zr at the interface between the eutectic structure and the matrix and at the grain boundaries. For the problem of reduced material plasticity, the eutectic phase is eliminated through homogenization treatment and regulation to prevent it from becoming a crack source. Taking all factors into consideration, it is necessary to control Zr in this application: 0.05% to 0.15%.

[0071] Mg: Mg has a high affinity for sulfur and readily forms sulfides, thus reducing the harmful effects of sulfur. Furthermore, Mg strongly segregates at grain and phase boundaries, where it refines carbides, reduces grain and phase boundary energies, enhances interfacial bonding, improves the alloy's creep strength and ductility, and hinders crack nucleation and propagation. For these reasons, the Mg content in this application should be controlled to ≤ 0.005%.

[0072] O: Reducing the O and N content can reduce the number of inclusions in the material, which helps improve the alloy's plasticity and toughness. During alloy melting, N easily reacts with Ti to form Ti(C, N). The increase in Ti(C, N) increases the potential for fatigue formation. Therefore, in this application, O: ≤ 20ppm and N: ≤ 20ppm are required.

[0073] S: A higher S content affects the plasticity and long-term performance of the alloy. At the end of alloy smelting and solidification, the higher the S content, the easier it is for sulfides to precipitate. The S element has a significant effect on nickel-based alloys above 800°C, and is more obvious during the smelting and billeting process of steel ingots. The inventors found in their research that the smelting of the experimental group with 100ppm S content failed, while the steel ingots with 56ppm S content cracked severely during the billeting process. When the S content exceeds 10ppm, the endurance life and plasticity at 730°C decrease to varying degrees. Therefore, it is necessary to control S: ≤10PPm in this application.

[0074] In order to further improve the comprehensive performance of the above-mentioned nickel-based high-temperature alloy, the components of the above-mentioned nickel-based high-temperature alloy include, by mass percentage: C: 0.04%~0.06%, Cr: 18.5%~20.0%, Co: 13.0%~14.0%, Mo: 4.0%~4.8%, Al: 1.3%~1.5%, Ti: 3.0%~3.25%, B: 0.015%~0.03%, Zr: 0.05%~0.15%, Ti / Al: 2.3~2.38, (Al+Ti): 4.4%~4.58%, Fe: 0.5%~1.0%, Mg≤0.005%, P: 0.004%~0.009%, O: ≤20PPm, N: ≤20PPm, S≤10PPm, nickel: balance.

[0075] Specifically, the microstructure of the nickel-based high-temperature alloy mainly includes equiaxed austenite grains and uniformly dispersed borides and carbides, as well as γ' phase dispersed inside the grains; wherein the carbides mainly include M 23 C6 and MC; among them M 23The C6 content is 0.85% to 1.17%, the MC content is 0.11% to 0.33%, and the boride content is 0.35% to 0.5%. The size of the γ' phase is small, and the size of the γ' phase is about 30 to 50 nm. The size of the MC carbide is about 5 to 10 μm. 23 The C6 size is about 0.3 to 1 μm.

[0076] Specifically, the M in the microstructure of the nickel-based high-temperature alloy 23 C6 is distributed in a continuous, necklace-like, or short rod-like form, and does not precipitate in a film-like form; MC is mostly in the form of blocks with smooth boundaries.

[0077] Specifically, the grains of the nickel-based high-temperature alloy are uniform and fine, and the average grain size can reach a uniform level of about 8 to 9.

[0078] Specifically, the properties of the nickel-based high-temperature alloy obtained in step 5 above are as follows: Room temperature properties: tensile strength σ b ≥1300MPa (e.g. 1330-1360MPa); yield strength σ 0.2 ≥1050MPa (e.g., 1052-1080MPa); elongation after fracture δ5 ≥27% (e.g., 27%-28%); cross-sectional shrinkage ψ ≥35% (e.g., 35%-40%); 730℃ / 550MPa endurance performance: endurance time τ ≥70h (e.g., 72-88h); elongation after fracture δ5 ≥17.5% (e.g., 17.7%-20%); 815℃ / 295MPa endurance performance: endurance time τ ≥83h (e.g., 83-95h); elongation after fracture δ5 ≥26% (e.g., 26%-30%); low-cycle fatigue performance: 500℃ / strain control 0-0.7% / 0.33Hz, >5×10 4 Week (e.g., 57532 to 63645 weeks).

[0079] The present application also provides applications of the nickel-based high-temperature alloy, which can be used for rotating parts of aircraft engines in service at temperatures above 700°C, such as alloy disk-shaft integrated turbine disks.

[0080] Compared with the prior art, the preparation method of the nickel-based high-temperature alloy of the present application reduces the cracking tendency by precisely controlling the process parameters of each step, for example, the homogenization heat treatment includes two stages of insulation; after the second stage of insulation, the furnace is cooled to below 1000±10°C and then air-cooled after being taken out of the furnace; combined with precise control of other steps and the parameters of each step, deformation is avoided, the alloy grains are ensured to be uniform and fine, and the morphology and distribution of carbides in the alloy and the content and size of other second phases are ensured to be optimally matched, thereby achieving the purpose of improving the alloy strength, reducing the cracking tendency, and reducing the crack propagation rate, thereby ensuring the comprehensive performance of the nickel-based high-temperature alloy.

[0081] The preparation method of the nickel-based high-temperature alloy of the present application improves the solid solution strengthening effect and grain boundary strength of the alloy by precisely controlling the content of individual elements such as B, Zr, Fe, C, Cr, Co, Al, and Ti in the alloy; and ensures the optimal matching of the morphology and distribution of carbides and the content and size of other second phases in the alloy by synergistically controlling the values ​​of B, Zr, and C. Without reducing the strength of the alloy, the risk sources of cracks in the alloy are reduced, the carbides are uniformly distributed at the grain boundaries of the alloy, the grain boundary strength is increased, and the speed of dislocation slip is slowed, thereby achieving the purpose of improving the strength of the alloy, reducing the cracking tendency, and reducing the crack propagation rate, thereby ensuring the comprehensive performance of the alloy.

[0082] The properties of the nickel-based high-temperature alloy prepared by the preparation method of the present application are as follows: Room temperature properties: tensile strength σ b ≥1300MPa (e.g. 1330-1360MPa); yield strength σ 0.2 ≥1050MPa (e.g., 1052-1080MPa); elongation after fracture δ5 ≥27% (e.g., 27%-28%); cross-sectional shrinkage ψ ≥35% (e.g., 35%-40%); 730℃ / 550MPa endurance performance: endurance time τ ≥70h (e.g., 72-88h); elongation after fracture δ5 ≥17.5% (e.g., 17.7%-20%); 815℃ / 295MPa endurance performance: endurance time τ ≥83h (e.g., 83-95h); elongation after fracture δ5 ≥26% (e.g., 26%-30%); low-cycle fatigue performance: 500℃ / strain control 0-0.7% / 0.33Hz, >5×10 4 Week (e.g., 57532 to 63645 weeks).

[0083] Examples 1-5

[0084] The advantages of the method for preparing the nickel-based high-temperature alloy of the present application are demonstrated below with specific examples and comparative examples. Examples 1-5 of the present application provide a method for preparing the nickel-based high-temperature alloy.

[0085] The compositions of the nickel-based high-temperature alloys of Examples 1-5 are shown in Table 1 below.

[0086] The preparation method of the nickel-based high-temperature alloy of Examples 1-5 includes:

[0087] Embodiment 1:

[0088] Step 1: sequentially performing vacuum induction melting, electroslag remelting, and vacuum arc remelting to obtain an ingot;

[0089] Step 2: Homogenize the ingot by annealing: heating from 400°C to 1165°C over 11 hours, keeping it warm for 37 hours; then continue heating to 1200°C, keeping it warm for 38 hours, then cooling it to 1000±10°C and air cooling it out of the furnace;

[0090] Step 3: Prepare the bar billet by rapid forging + radial forging: The alloy bar billet kept at 1200°C is subjected to three upsetting, drawing, and radial forging cycles in sequence; wherein, the deformation of each upsetting cycle is 30%, and the deformation of each drawing cycle is 30%; after each upsetting and drawing cycle, the temperature is lowered by 40°C until the holding temperature drops to 1080°C, and then the billet is forged in multiple passes on a radial forging press to a Φ180mm finished bar; the bar billet is forged using a die forging press to prepare forgings: The alloy bar billet segments kept at 1060°C are subjected to upsetting and die forging in sequence; the deformation of the upsetting and die forging cycles are 50% and 35%, respectively, to obtain forgings;

[0091] Step 4: Solution treatment: start heating from 400℃ at a rate of 5℃ / min to 1015℃, keep warm for 3.5h, and oil cool.

[0092] Step 5: Stabilization and aging treatment to obtain nickel-based high-temperature alloy discs: start heating from a furnace temperature of 400°C at a heating rate of 5°C / min, heat to 850°C and keep warm for 5 hours, air cool to 500±10°C, continue heating to 760°C and keep warm for 8 hours, and then air cool to room temperature to obtain nickel-based high-temperature alloy discs.

[0093] Example 2

[0094] The preparation method of this embodiment is roughly the same as that of Example 1, except that:

[0095] Step 1: performing vacuum induction melting and vacuum arc remelting in sequence to obtain an ingot;

[0096] Step 3: The alloy rod billet segment kept at 1060°C is subjected to upsetting, punching, and ring rolling in sequence; the deformation amounts of upsetting and ring rolling are 70% and 40% respectively, to obtain a forging;

[0097] Step 4: Air cooling is used for cooling the forging after solution treatment;

[0098] The rest of the steps are the same.

[0099] Example 3

[0100] The preparation method of this embodiment is roughly the same as that of Example 1, except that:

[0101] Step 1: sequentially performing vacuum induction melting, electroslag remelting, and vacuum arc remelting to obtain an ingot;

[0102] Step 3: Prepare a bar billet using a rapid forging method: The alloy bar billet, which is kept at 1180°C, is subjected to three upsetting, drawing, and radial forging cycles in sequence; wherein the deformation of each upsetting cycle is 50%, and the deformation of each drawing cycle is 50%; after each upsetting and drawing cycle, the temperature is lowered by 40°C until the holding temperature drops to 1060°C, after which the billet is cut into Φ240mm finished bars; the bar billet is forged using a die forging press to produce forgings: The alloy bar billet segments, which are kept at 1060°C, are subjected to upsetting and die forging in sequence; the deformation of the upsetting and die forging cycles are 50% and 35%, respectively, to obtain forgings;

[0103] The rest of the steps are the same.

[0104] The difference between Example 4-5 and Example 3 is that the chemical composition of the alloy is different, see Table 1, and the other parameters are the same.

[0105] This application also provides 6 comparative examples. The chemical compositions of the steels of Examples 1-5 and Comparative Examples 1-5 are shown in Table 1.

[0106] Comparative Examples 1-4 have different components from Example 1, but the preparation methods are the same.

[0107] The components of Comparative Example 5 are the same as those of Example 1, but the homogenization heat treatment method in the preparation method is different. The homogenization method in Comparative Example 5 is to heat the mixture from 400°C to 1160°C over 13 hours and keep it warm for 48 hours; then continue to heat the mixture to 1190°C, keep it warm for 66 hours, and then air cool it.

[0108] Comparative Example 6 has the same components as Example 1, but a different homogenization heat treatment method in the preparation method. In Comparative Example 6, the homogenization method is to keep the heat for 38 hours and then take the product out of the furnace for air cooling. Cracks appeared during the subsequent forging process.

[0109] Table 1 Chemical composition wt%

[0110] The metallographic structures of the embodiments and comparative examples are shown in Table 2.

[0111] Table 2 Metallographic structure

[0112] Table 3 shows the room temperature mechanical properties of the examples and comparative examples of the present application, Table 4 shows the endurance performance of the examples and comparative examples of the present application, and Table 5 shows the low cycle fatigue performance of the examples and comparative examples of the present application.

[0113] Table 3 Room temperature tensile mechanical properties

[0114] Table 4 Durability

[0115] Table 5 Low cycle fatigue properties at 500℃ / strain control 0~0.7% / 0.33Hz

[0116] Table 6 da / dN values ​​under different △K (stress intensity factor)

[0117] Figure 1 is a crack growth rate diagram of Example 1 and Comparative Example 4. It can be seen from the figure that under the same ΔK (stress intensity factor), the da / dN value (fatigue crack growth rate) corresponding to the present method is smaller. It can be seen that the crack growth rate of the nickel-based high-temperature alloy prepared by the preparation method of the present application is low.

[0118] Figures 2 and 3 are both microstructure diagrams of Example 1. As shown in Figure 2, blocky MC with smooth boundaries is distributed in the crystal, and as shown in Figure 3, short rod-shaped M 23 C6 is continuously and evenly distributed on the grain boundaries. Figure 4 is a microstructure diagram of comparative example 4, in which film-like M is distributed on the grain boundaries. 23 C6, FIG5 is a grain map of comparative example 4, and it can be seen that the grain structure of comparative example 4 is mixed crystal.

[0119] The above is only a preferred specific implementation method of the present application, but the scope of protection of the present application is not limited thereto. Any changes or replacements that can be easily thought of by any technician familiar with this technical field within the technical scope disclosed in this application should be covered by the scope of protection of the present application.

Claims

1. A preparation method of a nickel-based superalloy, characterized in that, Including: Step 1: Smelting to obtain an ingot; Step 2: Conducting homogenization heat treatment on the ingot; the homogenization heat treatment includes two-stage heat preservation; after the second-stage heat preservation, the furnace is cooled to below 1000 ± 10 °C and then taken out of the furnace for air cooling; Step 3: Forging to prepare bar blanks and forgings; Step 4: Solution treatment; Step 5: Stabilization and aging treatment to obtain a nickel-based superalloy.

2. The preparation method according to claim 1, characterized in that, In the said Step 2, the technological steps of the homogenization heat treatment include: S201: Slowly heating from a furnace temperature ≤ 400 °C to 1165 - 1170 °C, with a heating time of 10 - 15 h and a heat preservation time of 35 - 37 h; S202: Continuing to heat up to 1200 - 1210 °C, heat preserving for 38 - 42 h, then cooling the furnace to below 1000 ± 10 °C and taking it out of the furnace for air cooling.

3. The preparation method according to claim 1, characterized in that In the said Step 4, the solution treatment process is: heating to 995 - 1020 °C, heat preserving for 3 - 5 h, and taking it out of the furnace for cooling.

4. The preparation method according to claim 3, characterized in that, In the said Step 4, the heating rate is 3 - 6 °C / min.

5. The preparation method according to claim 3, wherein, In the said Step 4, for disk and shaft parts with a wall thickness of 20 - 200 mm, cooling is carried out by taking them out of the furnace and entering the cooling oil; for ring-shaped parts with a wall thickness of 10 - 100 mm, air cooling is used after taking them out of the furnace.

6. The preparation method according to claim 1, characterized in that, In the said Step 5, the stabilization and aging treatment process: heating to 820 - 850 °C, heat preserving for 5 - 7 h, air cooling to 500 ± 10 °C, continuing to heat up to 740 - 770 °C, heat preserving for 8 - 10 h, and then air cooling to room temperature.

7. The preparation method according to claim 6, characterized in that, In the said Step 5, the heating rate is controlled at 2 - 3.5 °C / min.

8. The preparation method according to any one of claims 1 to 7, characterized in that, The components of the said nickel-based superalloy include, by mass percentage: C: 0.02% - 0.06%, Cr: 18.5% - 20.0%, Co: 13.0% - 14.0%, Mo: 4.0% - 4.90%, Al: 1.3% - 1.6%, Ti: 2.80% - 3.25%, B: 0.015% - 0.03%, Zr: 0.05% - 0.15%, Ti / Al: 2.25 - 2.38, (Al + Ti): 4.35% - 4.58%, Fe: 0.05% - 2.0%, Mg ≤ 0.005%, P: 0.004% - 0.009%, O: ≤ 20 PPm, N: ≤ 20 PPm, S ≤ 10 PPm, with the balance being nickel.

9. The preparation method according to claim 8, characterized in that, The microstructure of the nickel-based superalloy obtained in the said Step 5 mainly includes equiaxed austenite grains, borides and carbides uniformly and dispersedly distributed, and γ' phases dispersedly distributed inside the grains.

10. The preparation method according to claim 9, characterized in that, In the microstructure of the nickel-based superalloy, the carbides mainly include M 23 C6 and MC.

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