A method for improving the stress-rupture resistance of powder nickel-base superalloys and articles made therefrom

Abstract

The application belongs to the field of high-temperature alloy, and relates to a preparation method for improving the sustained load crack propagation resistance of powder nickel-based high-temperature alloy and parts thereof. According to the final composition of the alloy, two component alloys with a difference of more than 25 DEG C in the gamma prime phase solid solution temperature are selected, the powders of the two component alloys are uniformly mixed, and then hot isostatic pressing or other thermal mechanical processing methods are used to prepare the alloy and part blanks, and a heat treatment system is selected according to the gamma prime phase solid solution temperature of the two component alloys. The powder nickel-based high-temperature alloy and parts prepared by the method can coordinate and improve the sustained load crack propagation resistance of the material by adjusting the composition and proportion of the two component alloys, and the service performance and life of the prepared aero-engine turbine disc are improved.

Description

Technical Field

[0001] This invention belongs to the field of powder superalloy technology and relates to a preparation method for improving the resistance of powder nickel-based superalloys and their components to load crack propagation, which is particularly suitable for the preparation of high-temperature rotating components in aero engines or gas turbines. Background Technology

[0002] High-temperature rotating components, especially high-pressure compressor disks and high-pressure turbine disks, are key components in aero engines. Their service performance largely determines the engine's performance, efficiency, and safety. Therefore, the high-temperature structural materials used must have higher operating temperatures and excellent comprehensive service performance, including high-temperature strength, oxidation resistance, creep resistance, and resistance to load-bearing crack propagation.

[0003] Nickel-based superalloys are key materials used in the manufacture of high-pressure compressors and high-pressure turbine rotating components in aero-engines. They are primarily produced as blanks through casting-forging or powder metallurgy processes, and then machined into final parts. Because powder metallurgy produces materials and parts with a uniform and fine microstructure, they exhibit better service performance than those produced by casting-forging. Furthermore, powder metallurgy can overcome macroscopic element segregation, making it more suitable for the preparation of highly alloyed, high-temperature-resistance superalloy materials and parts.

[0004] Currently, rotating parts made from powder metallurgy high-temperature alloys are produced by processing alloy powders prepared from single-component alloy ingots through argon atomization or rotating electrode methods, followed by thermomechanical and heat treatment processes such as hot isostatic pressing, extrusion, and forging. These materials exhibit excellent compositional and microstructural uniformity. For example... Figure 1 As shown, for powder superalloys and their components prepared by this existing route, when heat-treating the final part blank, solution heat treatment can be performed above or below the solution temperature of the γ' strengthening phase to obtain a uniform over-solution coarse-grained structure or a sub-solution fine-grained structure. For uniformly structured materials prepared by existing processes, the performance control of the alloy and components is mainly achieved by adjusting the microstructure (including grains, γ' strengthening phase, etc.) in the final product through alloy design, extrusion, forging, and heat treatment processes during preparation to obtain the required performance.

[0005] To improve the heat resistance of powdered nickel-based superalloys and enhance their high-temperature strength, creep resistance, and fatigue resistance, more alloying elements are added to increase the volume fraction of the γ' strengthening phase. Currently, most powdered nickel-based superalloys with high heat resistance have a γ' strengthening phase volume fraction ≥45% at room temperature. This method reduces the material's resistance to brittle intergranular fracture caused by oxidation at high temperatures, resulting in poor resistance to load-holding crack propagation and significantly impacting the damage tolerance of components. To improve the load-holding crack propagation resistance of powdered nickel-based superalloys with high γ' strengthening phase content, one approach is to reduce the cooling rate of the material or component blank after solution heat treatment. However, this method sacrifices the material's tensile strength and creep resistance, and its effect on improving load-holding crack propagation resistance is limited. Another method is to form a necklace-like structure in the material through thermomechanical deformation. However, this method is difficult to control in practical engineering, making it hard to obtain a controllable and uniformly distributed beneficial structure.

[0006] In summary, it is currently difficult to coordinate and improve the key performance required for service, such as the resistance to crack propagation under load, high-temperature strength, and creep resistance, of high-temperature nickel-based superalloys with high temperature resistance. Summary of the Invention

[0007] The purpose of this invention is to design a preparation method for high-temperature rotating components in aero-engines or gas turbines, such as high-pressure compressor disks and high-pressure turbine disks in aero-engines, especially powdered nickel-based superalloy materials and their components with a γ' strengthening phase volume fraction ≥45% at room temperature, which can coordinately improve key performance required for service, such as resistance to load-bearing crack propagation, high-temperature strength, and creep resistance.

[0008] To solve this technical problem, the technical solution of the present invention is as follows:

[0009] A method for improving the resistance to load-bearing crack propagation of powdered nickel-based superalloys and their components, wherein the method designs two alloys with different compositions as component alloys based on the chemical composition of the target alloy, and the composition of the mixture of the two component alloys in a certain mass ratio is the same as that of the target alloy; the difference in solid solution temperature of the γ' phase of the two component alloys needs to be ≥25℃.

[0010] The two alloy components are uniformly mixed according to the mass ratio, and a part blank is prepared by thermomechanical processing; the microstructure of the part blank is controlled by heat treatment.

[0011] The solution heat treatment temperature during the heat treatment process is selected based on the solution temperature of the γ' phase in the two-component alloy as follows:

[0012] The solution heat treatment temperature is between the solution temperatures of the γ' phase in the two-component alloy, resulting in a dual-distribution grain structure of subsolid solution and oversolid solution.

[0013] The solution heat treatment temperature is simultaneously higher than the solution temperature of the γ' phase in both alloy components, resulting in a dual-distribution grain structure of over-solution + over-solution. The solution heat treatment temperature is 0 to 30°C higher than the higher solution temperature of the γ' phase in the two alloy components.

[0014] The preparation method steps are as follows:

[0015] Step 1: Based on the chemical composition of the target alloy, design two component alloys A and B with different compositions;

[0016] Step 2: Prepare the raw materials for nickel-based superalloys A and B separately, and prepare the master alloys by vacuum induction melting respectively;

[0017] Step 3: Prepare alloy powders from the two types of master alloys and sieve them to the selected particle size.

[0018] Step 4: Mix the two alloy powders after sieving in a certain proportion.

[0019] Step 5: Load the mixed alloy powder into a hot isostatic pressing (HIP) sleeve, and then degas and seal it.

[0020] Step 6: Place the hot isostatic pressing (HIP) package into a hot isostatic pressing (HIP) apparatus for densification treatment to obtain a hot isostatic pressing ingot or part blank; if the obtained ingot is a hot isostatic pressing (HIP) ingot, perform forming thermomechanical treatments such as extrusion or forging to obtain a part blank.

[0021] Step 7: Perform heat treatment on the part blank obtained in Step 6; wherein the solution heat treatment temperature during the heat treatment process is selected according to the solution temperature of the γ' phase in the two component alloys A and B.

[0022] Step 8: Machine the blank part into a part.

[0023] The thermomechanical processing process can be one or a combination of several methods such as hot isostatic pressing, hot sintering, extrusion, forging, and additive manufacturing.

[0024] The target composition of the nickel-based powder superalloy is as follows: Cr 5.0%–20.0%, Co 5.0%–25.0%, W 0.0%–15.0%, Mo 0.0%–5.0%, Al 0.0%–6.0%, Ti 0.0%–6.0%, Nb 0.0%–8.0%, Ta 0.0%–6.0%, Hf 0.0%–1.0%, Zr 0.0%–0.2%, C 0.01%–0.3%, B 0.01%–0.3%, with the balance being Ni, other trace elements beneficial to superalloys, and unavoidable trace or impurity elements.

[0025] Preferably, the target composition of the nickel-based powder superalloy is as follows: Cr 7.0%–18.0%, Co 7.0%–23.0%, W 1.0%–12.0%, Mo 1.0%–3.5%, Al 2.0%–5.0%, Ti 0.5%–5.0%, Nb 0.5%–8.0%, Ta 0.0%–6.0%, Hf 0.0%–0.5%, Zr 0.0%–0.2%, C 0.01%–0.2%, B 0.01%–0.2%, with the balance being Ni, other trace elements beneficial to the superalloy, and unavoidable trace or impurity elements.

[0026] Furthermore, the composition is as follows: Cr 8.0%–16.0%, Co 8.0%–19.0%, W 2.0%–11.0%, Mo 1.4%–3.2%, Al 2.1%–5.0%, Ti 0.5%–4.3%, Nb 0.5%–7.9%, Ta 0.0%–5.0%, Hf 0.0%–0.4%, Zr 0.0%–0.1%, C 0.02%–0.15%, and B 0.015%–0.025%.

[0027] The preparation method of this invention is used for the preparation of high-temperature rotating parts in aero-engines or gas turbines, and is applicable to the preparation of nickel-based powder superalloys, especially applicable to the final alloy material in which the volume fraction of the γ' strengthening phase is ≥45% at room temperature.

[0028] The beneficial effects of this invention are:

[0029] The preparation method of this invention, through design or selection, prepares materials with two γ' phases having different solution temperatures, allowing for different grain size changes in the two components during subsequent solution heat treatment. If the solution heat treatment temperature is between the solution temperatures of the γ' phases of the two component alloys, a material microstructure with a uniformly distributed mixture of small-sized subsolidified grains and large-sized oversolidified grains can be obtained. If the solution temperature is simultaneously higher than the solution temperatures of the γ' phases of both component alloys, although oversolidified grain growth occurs in both components, the different initial chemical compositions result in different grain growth rates, thus still yielding a dual-distribution oversolidified grain microstructure. Furthermore, regardless of the heat treatment method, the distribution of the γ' phase within the grains differs simultaneously with the different grain growth processes.

[0030] Therefore, due to differences in grain size and intragranular γ' phase distribution, the present invention actually yields a material with non-uniformly distributed properties at the meso-microscale. The resulting material microstructure exhibits localized non-uniform deformation and stress distribution under load. This non-uniform deformation and stress distribution can reduce the stress-strain levels at certain grain boundaries, thereby reducing the driving force for load-holding crack propagation and consequently lowering the load-holding crack propagation rate in the material. Furthermore, due to mechanisms such as synergistic strengthening and synergistic deformation, this microstructure type is even beneficial to the tensile strength and plasticity of the material. Since solution heat treatment does not require reducing the cooling rate to improve the load-holding crack propagation resistance, the method of the present invention also helps materials and parts maintain tensile strength and creep resistance. Attached Figure Description

[0031] Figure 1 This is a schematic diagram comparing the process flow of the present invention with that of existing processes;

[0032] Figure 2 These are schematic diagrams showing the microstructures of the alloys prepared by the process flow of the present invention and the existing process flow in Examples 1, 2, 3, and 4. Figure 2 (a) shows the microstructure of a single-component alloy A after subsolution heat treatment obtained using the heat treatment temperatures in each embodiment, where hexagons represent grain structure and black dots at the hexagonal boundaries represent the primary γ' phase in the material after subsolution heat treatment. Figure 2 (b) shows the microstructure of a single-component B alloy after solution heat treatment obtained using the heat treatment temperatures in each embodiment, where hexagons represent grain structures (no primary γ' phase at the boundaries). Figure 2 (c) Alloy C is prepared by solution heat treatment using the process of the present invention at temperatures lower than the solution temperature of the γ' phase of alloy A and higher than the solution temperature of the γ' phase of alloy B in each embodiment. The large hexagon represents the over-solution large-sized grains evolved from alloy B, and the small hexagon represents the sub-solution small-sized grains evolved from alloy A (the black dots at the interface are still the primary γ' phase). Figure 2 (d) shows the microstructure of the comparative alloys obtained by subsolution heat treatment using a single component C and the heat treatment temperatures in each embodiment according to the existing process; it can be seen from the figure that the alloys prepared by the process of the present invention have a dual-distributed grain structure, which is significantly different from the uniform material structure prepared by the existing process.

[0033] Figure 3 This is a schematic diagram of the microstructure of the alloy prepared by the process flow of the present invention and the existing process flow in Example 5, wherein, Figure 3 (a) shows the microstructure of a single-component alloy A after solution heat treatment, obtained using the heat treatment temperatures in each embodiment, where hexagons represent grain structures. Figure 3(b) The microstructure of the single-component B alloy obtained by solution heat treatment using the heat treatment temperatures in each embodiment. Figure 3 (c) Alloy C is prepared by solution heat treatment at temperatures higher than the γ' phase solution temperature of both alloy A and alloy B according to the process of the present invention in each embodiment. The large hexagon represents the larger oversolution grains evolved from alloy B, and the small hexagon represents the smaller oversolution grains evolved from alloy A. Figure 3 (d) shows the microstructure of the oversolution heat treatment obtained by using a single component C and the heat treatment temperature in each embodiment according to the existing process; it can be seen from the figure that the alloy prepared by the process of the present invention has a dual-distributed oversolution grain structure, which is significantly different from the uniform oversolution material structure prepared by the existing process. Detailed Implementation

[0034] To make the objectives, technical solutions, and advantages of the embodiments of the present invention clearer, the technical solutions of the embodiments of the present invention will be clearly and completely described below with reference to the accompanying drawings. Obviously, the described embodiments are only some, not all, of the embodiments of the present invention. Based on the embodiments of the present invention, all other embodiments obtained by those skilled in the art without creative effort are within the scope of protection of the present invention.

[0035] The features of various aspects of the embodiments of the present invention will now be described in detail. In the following detailed description, numerous specific details are set forth in order to provide a thorough understanding of the invention. However, it will be apparent to those skilled in the art that the invention can also be practiced without these specific details. The following description of the embodiments is merely intended to provide a better understanding of the invention by illustrating examples. The invention is not limited to any specific setups and methods provided below, but covers all improvements, substitutions, etc., to product structures and methods without departing from the spirit of the invention.

[0036] In the various accompanying drawings and the following description, well-known structures and techniques are not shown in order to avoid unnecessarily obscuring the invention.

[0037] Example 1:

[0038] This embodiment consists of Comparative Example 1 and Examples 1-1 to 1-4 (material composition details are shown in Table 1). Examples 1-1 to 1-4 were prepared by mixing component A and component B according to the proportions in the table using the process described in this invention. The proportions of component A and component B include 95%:5%, 90%:10%, 80%:20%, and 50%:50% (all are mass fractions).

[0039] The comparative examples have the same equivalent composition as the mixture of components A and B in the proportions shown in the table in each example, and are prepared using existing processes (i.e., no powder mixing step, resulting in a uniform grain structure).

[0040] In all the examples in this series (including the comparative examples), the final solution heat treatment temperature was 1140°C. This solution heat treatment temperature is between the solution temperatures of the γ' phase of components A and B in Examples 1-1 to 1-4, and lower than the solution temperature of the γ' phase of the materials in the comparative examples. Therefore, Comparative Example 1 yielded a uniform subsolid-solid grain structure, while Examples 1-1 to 1-4 yielded a subsolid-solid grain structure with coarser oversolid-solid grains distributed within it.

[0041] The following description will be based on Example 1-1:

[0042] A powdered nickel-based superalloy prepared from a two-component alloy powder, the target composition of which is shown in Examples 1-1 of Table 1. The composition of this alloy is equivalent to 95% component A alloy plus 5% component B alloy in Examples 1-1 of Table 1, and is the same as Comparative Example 1 (i.e., composition C).

[0043] First, vacuum induction melting was used to prepare the elemental materials into component A and component B master alloy ingots corresponding to Examples 1-1 in Table 1. Then, argon atomization was used to prepare alloy powders of component A and component B, respectively, and the powders were sieved to -270 mesh. The solid solution temperature of the γ' strengthening phase in component A alloy is approximately 1160℃, and the solid solution temperature of the γ' strengthening phase in component B alloy is approximately 1120℃.

[0044] 95% by mass of component A alloy powder and 5% by mass of component B alloy powder were uniformly mixed, then placed in a steel sleeve, and subjected to degassing, sealing welding, and hot isostatic pressing. The dense billet after hot isostatic pressing was then subjected to solution heat treatment at 1140℃ (a temperature between the solution temperatures of the two component alloys). This process caused component B to undergo solution heat treatment, thus transforming it into a coarse-grained structure; while component A underwent a semi-solution heat treatment, retaining a fine-grained structure. Figure 1 As shown in the schematic diagram, the alloy obtained by this method has two grain size distributions: a small number of large grains with a size of about 20 to 40 μm are uniformly distributed in fine grains with a size of about 10 μm, and primary γ' strengthening phases with a size of about 1 μm are also dispersed at the grain boundaries of the small grains.

[0045] Comparative Example 1, prepared using existing processes, involves using only the target alloy component C to prepare the master alloy and alloy powder, followed by the same hot isostatic pressing and heat treatment processes. Since the solution temperature of the γ' strengthening phase in Comparative Example 1 alloy is 1156℃, higher than the solution heat treatment temperature of 1140℃, the resulting alloy has a uniform subsolid-solid-solid grain structure. The difference between this and the alloy in this case is the absence of dispersed, uniformly distributed large grains. Compared to the target alloy prepared using this existing process, the bi-grain size distribution alloy prepared using the bi-component alloy mixing method of this invention exhibits a significantly reduced crack propagation rate under load. For example, under test conditions of 650℃ and a maximum stress of 2 minutes, the crack propagation rate at ΔK = 30 MPa decreased to approximately 60%; simultaneously, the material's room temperature tensile properties, tensile properties at 650℃, and creep properties at 650℃ remained unchanged.

[0046] Examples 1-2 to 1-4 were prepared using the same process as Example 1-1, with only the ratio of component A and component B being changed. Compared to Comparative Example 1, Examples 1-2 to 1-4 also exhibited better resistance to load-maintaining crack propagation.

[0047] Examples 1-2 to 1-4 can also be prepared by hot isostatic pressing (HIP) + extrusion + forging. After HIP, the mixed powder is hot-extruded and isothermally forged at a billet temperature simultaneously lower than the solution temperatures of both alloys to obtain a part blank with a specific shape. Then, a solution heat treatment at 1140°C is performed to obtain an alloy material with a bicrystalline grain size distribution. Its resistance to crack propagation under load is comparable to that of the material obtained by direct HIP, and both are superior to the alloy in Comparative Example 1 prepared using the same process.

[0048] Table 1

[0049]

[0050]

[0051] Example 2:

[0052] This embodiment includes Comparative Example 2 and Examples 2-1 to 2-3 (material composition details are shown in Table 2). This embodiment is similar to Example 1, but the chemical composition of the comparative examples and examples differs from that of Example 1. In this embodiment, the ratio of component A to component B includes 90%:10%, 80%:20%, and 60%:40% (all are mass fractions).

[0053] In Example 2, the heat treatment temperature of the comparative example and the embodiment was 1173°C, which is lower than that of the comparative example and similarly falls between the solid solution temperatures of the γ' phase of components A and B, and is lower than the solid solution temperature of the γ' phase of the material in the comparative example. Compared to the comparative example, when the above embodiments were held at the highest stress of 650–700°C for 2 minutes, the crack propagation rate decreased to approximately 10–60% at ΔK = 30 MPa.

[0054] Table 2

[0055]

[0056] Example 3:

[0057] This embodiment includes Comparative Example 3 and Examples 3-1 and 3-2 (material composition details are shown in Table 3). This embodiment is similar to Examples 1 and 2, but the chemical composition of the comparative examples and embodiments differs from that of Examples 1 and 2. In this embodiment, the ratio of component A to component B includes 95%:5% and 90%:10% (both are mass fractions).

[0058] In Example 3, the heat treatment temperature of the comparative example and the embodiment is 1165°C, which is lower than that of the comparative example and similarly between the solid solution temperatures of the γ' phase of component A and component B, and lower than the solid solution temperature of the γ' phase of the material in the comparative example.

[0059] Table 3

[0060]

[0061] Example 4:

[0062] This embodiment includes four sets of comparisons with different chemical compositions (see Table 4 for material composition details). In the embodiments, the ratio of component A to component B is 90%:10% (mass fraction). The heat treatment temperatures of the comparative examples and embodiments in Examples 4-1 to 4-4 are all between the γ' phase solid solution temperatures of component A and component B, and lower than the γ' phase solid solution temperatures of the materials in the comparative examples.

[0063] Table 4

[0064]

[0065]

[0066] Example 5:

[0067] This embodiment includes four sets of comparisons with different chemical compositions. Each set of chemical compositions is the same as that in Example 4, and the material compositions are detailed in Table 5. In this embodiment, the ratio of component A to component B is 90%:10% (mass fraction). The heat treatment temperatures of the comparative examples and embodiments in Examples 5-1 to 5-4 are all simultaneously higher than the γ' phase solid solution temperatures of components A and B, as well as the γ' phase solid solution temperature of the material in the comparative examples. Compared with the comparative examples, the crack propagation rate of the above embodiments is reduced to approximately 5% to 50% when held at the highest stress of 650–750°C for 2 minutes at ΔK = 30 MPa.

[0068] Table 5

[0069]

[0070] In the embodiments obtained by this process, both component A and component B evolved into coarse-grained grains after solution heat treatment. However, due to the different initial chemical compositions, the grain sizes and intragranular strengths after heat treatment were different.

[0071] Finally, it should be noted that the above embodiments are only used to illustrate the technical solutions of the present invention, but the protection scope of the present invention is not limited thereto. Any person skilled in the art can easily conceive of various equivalent modifications or substitutions within the technical scope disclosed in the present invention, and these modifications or substitutions should be covered within the protection scope of the present invention.

Claims

1. A method for preparing powdered nickel-based superalloys to improve the resistance to crack propagation under load, characterized in that: The preparation method involves designing two alloys with different compositions as component alloys based on the chemical composition of the target alloy. The composition of the two component alloys after being mixed in a certain mass ratio is the same as that of the target alloy. The solid solution temperature difference of the γ' phase between the two component alloys needs to be ≥ 25℃. The two alloy components are uniformly mixed according to the mass ratio, and a part blank is prepared by thermomechanical processing; the microstructure of the part blank is controlled by heat treatment. The solution heat treatment temperature during the heat treatment process is selected based on the solution temperature of the γ' phase in the two-component alloy as follows: The solution heat treatment temperature is between the solution temperatures of the γ' phase in the two-component alloy, resulting in a dual-distribution grain structure of subsolid solution and oversolid solution. The solution heat treatment temperature is simultaneously higher than the solution temperature of the γ' phase in the alloy of the two components, resulting in a dual-distribution grain structure of over-solution + over-solution. The solution heat treatment temperature is 0~30℃ higher than the higher solution temperature of the γ' phase in the alloy of the two components. The target composition of the powdered nickel-based superalloy is as follows: Cr 5.0%~20.0%, Co 5.0%~25.0%, W 0.0%~15.0%, Mo 0.0%~5.0%, Al 0.0%~6.0%, Ti 0.0%~6.0%, Nb 0.0%~8.0%, Ta 0.0%~6.0%, Hf 0.0%~1.0%, Zr 0.0%~0.2%, C 0.01%~0.3%, B 0.01%~0.3%, with the balance being Ni and unavoidable trace or impurity elements.

2. The preparation method according to claim 1, characterized in that: The preparation method steps are as follows: Step 1: Based on the chemical composition of the target alloy, design two component alloys A and B with different compositions; Step 2: Prepare the raw materials for nickel-based superalloys A and B separately, and prepare the master alloys by vacuum induction melting respectively; Step 3: Prepare alloy powders from the two types of master alloys and sieve them to the selected particle size. Step 4: Mix the two alloy powders after sieving in a certain proportion. Step 5: Load the mixed alloy powder into a hot isostatic pressing (HIP) sleeve, and then degas and seal it. Step 6: Place the hot isostatic pressing (HIP) jacket inside the HIP equipment for densification treatment to obtain a hot isostatic pressing ingot or part blank; if the obtained ingot is a hot isostatic pressing ingot, then perform extrusion, forging, and thermomechanical treatment to obtain a part blank. Step 7: Perform heat treatment on the part blank obtained in Step 6; wherein the solution heat treatment temperature during the heat treatment process is selected according to the solution temperature of the γ' phase in the two component alloys A and B; Step 8: Machine the blank part into a part.

3. The preparation method according to claim 1, characterized in that: The thermomechanical processing method is one or a combination of several of the following methods: hot isostatic pressing, hot sintering, extrusion, forging, and additive manufacturing.

4. The preparation method according to claim 1, characterized in that: The target composition of the powdered nickel-based superalloy is as follows: Cr 7.0%~18.0%, Co 7.0%~23.0%, W 1.0%~12.0%, Mo 1.0%~3.5%, Al 2.0%~5.0%, Ti 0.5%~5.0%, Nb 0.5%~8.0%, Ta 0.0%~6.0%, Hf 0.0%~0.5%, Zr 0.0%~0.2%, C 0.01%~0.2%, B 0.01%~0.2%, with the balance being Ni and unavoidable trace or impurity elements.

5. The preparation method according to claim 4, characterized in that: Cr 8.0%~16.0%, Co 8.0%~19.0%, W2.0%~11.0%, Mo 1.4%~3.2%, Al 2.1%~5.0%, Ti 0.5%~4.3%, Nb 0.5%~7.9%, Ta 0.0%~5.0%, Hf 0.0%~0.4%, Zr 0.0%~0.1%, C 0.02%~0.15%, B 0.015%~0.025%.

6. The preparation method according to claim 1, characterized in that: The preparation method is applicable to the preparation of nickel-based powder superalloys, and the volume fraction of the γ' strengthening phase in the final alloy material is ≥45% at room temperature.

7. The preparation method according to claim 1, characterized in that: The preparation method is used for the preparation of high-temperature rotating components in aero engines or gas turbines.

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