Method for improving utilization rate of nickel-based powder superalloy and application

Abstract

The invention discloses a method for improving the utilization rate of nickel-based powder superalloy, which comprises the following steps: preparing alloy powder from mother alloy through a powder preparation process, screening the alloy powder to obtain powder with the particle size of less than 150 microns, and dividing the powder into fine powder and coarse powder; wherein the fine powder refers to powder with the particle size smaller than 53 micrometers, the coarse powder refers to powder not smaller than 53 micrometers, and the alloy is prepared from the fine powder and the coarse powder through a forming process. The invention further provides application of the alloy prepared through the method to hot end components of aerospace, ships and energy power systems. The applicant provides the method for improving the utilization rate of the nickel-based powder high-temperature alloy and the related alloy, and verifies that coarse powder with the particle size of 53-150 microns and fine powder with the particle size of 18-53 microns are on the same level in structure form and comprehensive mechanical property, and the powder utilization rate is greatly improved.

Description

Technical Field

[0001] This application relates to the field of powder metallurgy technology, and in particular to a method and application for improving the utilization rate of nickel-based powder superalloys. Background Technology

[0002] Powder metallurgy nickel-based superalloys have excellent high-temperature strength, creep and fatigue crack propagation resistance, and are widely used in hot-end components that need to withstand high temperatures, such as high-pressure turbine disks for aero engines.

[0003] The process of manufacturing turbine disks and other workpieces from nickel-based superalloys mainly involves multiple steps, including high-purity master alloy preparation, powder preparation, hot isostatic pressing, hot extrusion, isothermal forging, heat treatment, and machining. Raw material utilization is low, with a material utilization rate of less than 20% from raw materials to the finished turbine disk. Among these steps, the yield rate of powder preparation is the most significant factor leading to this substantial reduction in material utilization. The main purpose of collecting fine powder with smaller particle sizes is to reduce the impact of non-metallic inclusions on fatigue performance; that is, the finer the collected powder, the smaller the size of the non-metallic inclusions, which is beneficial for improving the fatigue performance of the alloy.

[0004] Currently, the main powder preparation methods include atomization powder preparation and rotating electrode powder preparation. Atomization powder preparation typically collects fine powder smaller than 53 micrometers for subsequent disk production, with a fine powder yield of approximately 60%. Rotating electrode powder preparation yields even lower fine powder compared to argon atomization powder preparation. Low powder yield is one of the main reasons for the high cost of powder metallurgy turbine disks. Literature data shows that collecting only powder smaller than 45-53 micrometers will inevitably further reduce the powder yield and increase production costs. Summary of the Invention

[0005] To address the aforementioned technical problems, the first objective of this invention is to provide a method for improving the utilization rate of nickel-based powder superalloys; the second objective of this invention is to provide applications of alloys prepared by the above method. In comparison with current powder preparation applications, the applicant has proposed a method for improving the utilization rate of nickel-based powder superalloys and related alloys, demonstrating that coarse powder with a particle size of 53-150 μm and fine powder with a particle size of 18-53 μm have similar microstructures and comprehensive mechanical properties, resulting in a significant improvement in powder utilization.

[0006] The technical solution provided by this invention is as follows: A method for improving the utilization rate of nickel-based powder superalloys involves obtaining alloy powder from a master alloy through a powder-making process, sieving the alloy powder to obtain powder with a particle size of less than 150 μm, and separating it into fine powder and coarse powder; wherein fine powder refers to powder with a particle size of less than 53 micrometers, and coarse powder refers to powder with a particle size of not less than 53 micrometers, and the fine powder and coarse powder are respectively used to prepare alloys through forming processes.

[0007] Preferably, the powder preparation process is one of vacuum atomization powder preparation, electrode induction gas atomization powder preparation, and rotating electrode powder preparation. Preferably, the powder preparation process uses high-purity argon gas as the atomization medium.

[0008] Preferably, the alloy powder is sieved to obtain fine powder with a particle size greater than or equal to 18 μm and less than 53 μm, and coarse powder with a particle size greater than or equal to 53 μm and less than 150 μm, and the target alloy is prepared by forming process respectively.

[0009] Preferably, the target alloy is prepared by forming the coarse powder. Preferably, the process includes the following steps: S1. Prepare the master alloy by vacuum induction melting process of each component; S2. The master alloy is processed into alloy powder by gas atomization, and the alloy powder is sieved to obtain powder. S3. The powder is processed by hot isostatic pressing. S4. Perform hot deformation treatment on the alloy after hot isostatic pressing. S5. Perform heat treatment on the alloy after hot deformation.

[0010] Preferably, step S3 further includes the steps of cladding preparation, cladding welding, powder loading, degassing, and sealing welding before hot isostatic pressing; After step S3 and before step S4, the process also includes removing the cladding and then using a stainless steel cladding to externally wrap the hot isostatic pressed alloy ingot. In step S4, the hot deformation treatment includes either hot extrusion or hot forging. In step S5, the heat treatment includes any one or both of solution heat treatment and aging heat treatment.

[0011] Preferably, in step S1, the casting temperature of vacuum induction melting is 1400℃-1550℃; In step S3, the hot isostatic pressing temperature is 1050℃-1200℃, the pressure is 100-180MPa, and the time is 3-8 hours; In step S4, the hot extrusion temperature is 1050-1150℃, and the extrusion ratio is 4-8; The hot forging temperature range is 1050-1150℃, and the rate is less than 0.01 / s; In step S5, the solution heat treatment is performed by heating to 1130-1200℃, holding at that temperature for 1-4 hours, and then rapidly cooling to room temperature. The aging heat treatment is as follows: heat to 750-800 ℃, hold for 6-12 hours, and then air cool to room temperature.

[0012] Preferably, the master alloy, by mass fraction, comprises the following components: Co 24%–30%; Cr 11%–14%; Mo 3.5%–5%; W 1%–2.8%; Al 2%–5%; Ti 2%–4%; Nb 1%–3%; Ta 2%–3.8%; C 0.02%–0.1%; B 0.01%–0.1%; Zr 0.02%–0.08%; the balance being Ni and impurities; preferably, the master alloy, by mass fraction, comprises the following components: Co 24%–26%; Cr 11%–13%; Mo 4%–5%; W 1.5%–2.5%; Al 2%–4%; Ti 2%–4%; Nb 1%–2%; Ta 3%–3.8%; C 0.03%–0.05%; B 0.02%–0.04%; Zr 0.04% to 0.06%; the balance is Ni and impurities.

[0013] More preferably, the master alloy comprises, by mass fraction, the following components: Co 25%–25.3%; Cr 11.8%–12.2%; Mo 3.8%–4.2%; W 1.8%–2.2%; Al 2.8%–3.2%; Ti 2.7%–3.3%; Nb 1.3%–1.7%; Ta 3.2%–3.7%; C 0.03%–0.05%; B 0.02%–0.04%; Zr 0.04%–0.06%; with the balance being Ni and impurities.

[0014] More preferably, the master alloy comprises, by mass fraction, the following components: Co 25.0%; Cr 12.0%; Mo 4.0%; W 2.0%; Al 3.0%; Ti 3.0%; Nb 1.5%; Ta 3.5%; C 0.04%; B 0.03%; Zr 0.05%; with the balance being Ni and impurities.

[0015] In the preferred high-cobalt, low-tantalum-nickel-based alloy formulations, the amounts of cobalt and tantalum meet the requirements: 4.024%≤0.326Co-Ta≤7.78%, and the Co content ≥24% and the Ta content ≤3.8%.

[0016] The above high-cobalt, low-tantalum-nickel-based alloy formulation has a heat treatment window exceeding 100℃, which improves the alloy's temperature resistance, reduces the solution heat treatment temperature and cost, and lowers the risk of initial melting during solution heat treatment. Simultaneously, the alloy's resistance to creep deformation is also significantly enhanced.

[0017] Alternatively, the master alloy, by mass fraction, comprises the following components: Co 10%–14%; Cr 10%–14%; Mo 3.5%–5%; W 1%–2.8%; Al 2%–5%; Ti 2%–4%; Nb 0.5%–2.5%; Ta 4%–7%; C 0.02%–0.1%; B 0.02%–0.1%; Zr 0.02%–0.1%; the balance being Ni and impurities; preferably, the master alloy, by mass fraction, comprises the following components: Co 11%–13%; Cr 11%–13%; Mo 4%–5%; W 1.5%–2.5%; Al 2%–4%; Ti 2%–4%; Nb 0.5%–2.5%; Ta 4%–6%; C 0.03%–0.05%; B 0.02%–0.04%; Zr 0.04% to 0.06%; the balance is Ni and impurities.

[0018] Preferably, the master alloy comprises, by mass fraction, the following components: Co 12%; Cr 12%; Mo 4%; W 2%; Al 3%; Ti 3%; Nb 1.5%; Ta 5%; C 0.04%; B 0.03%; Zr 0.05%; with the balance being Ni and impurities.

[0019] The aforementioned low-cobalt nickel-based alloy formulation improves the alloy's creep resistance by reducing cobalt and increasing the amount of tantalum, thereby addressing the problem that existing nickel-based high-temperature alloy materials have difficulty improving solidus temperature and heat treatment window.

[0020] Alternatively, the master alloy, by mass fraction, comprises the following components: Co 10%–14%; Cr 10%–14%; Mo 4%–6%; W 0.15%; Al 2%–4%; Ti 2%–4%; Nb 0.75%–1.4%; Ta 6.2%–8.5%; C 0.11%–0.2%; B 0.02%–0.1%; Zr 0.02%–0.1%; the balance being Ni and impurities.

[0021] Preferably, the master alloy comprises, by mass fraction, the following components: Co 11%–13%; Cr 11%–13%; Mo 4%–6%; W 0–1.5%; Al 2%–4%; Ti 2%–4%; Nb 0.75%–1.25%; Ta 6.2%–7.8%; C 0.14%–0.16%; B 0.02%–0.04%; Zr 0.04%–0.06%; with the balance being Ni and impurities.

[0022] More preferably, the master alloy comprises, by mass fraction, the following components: Co 12%; Cr 12%; Mo 5%; W 0%; Al 3%; Ti 3%; Nb 1%; Ta 7%; C 0.15%; B 0.03%; Zr 0.05%; with the balance being Ni and impurities.

[0023] The above-mentioned high-tantalum, high-carbon nickel-based alloy formulation increases the carbon content, resulting in the precipitation of high-content carbides at the grain boundaries, which inhibits grain growth and improves the high-temperature strength of the alloy. At the same time, it also increases the solid solution temperature of the γ' precipitate phase and improves the service temperature of the alloy by reducing the cobalt content and increasing the tantalum content.

[0024] Alloys prepared by the method described in any of the above methods are used in the manufacture of hot-end components for aerospace, shipbuilding, and energy and power systems.

[0025] This application first provides a method for improving the utilization rate of nickel-based powder superalloys. The master alloy is processed into alloy powder through a powder making process. The alloy powder is then sieved to obtain powder with a particle size of less than 150 μm. The powder yield exceeds 90%, and the alloy is prepared through a forming process.

[0026] Preferably, the alloy powder is sieved to obtain powder with a particle size greater than or equal to 18 μm and less than 53 μm (referred to as fine powder) and powder with a particle size greater than or equal to 53 μm and less than 150 μm (referred to as coarse powder). The yield of coarse powder is about 30-40%, and the yield of fine powder is about 45-55%. The total yield of these two different particle size ranges of powder exceeds 85%.

[0027] The applicant conducted comparative experiments on powders with different particle sizes, selecting alloy powders with the same formula and particle sizes of 53-150 μm and 18-53 μm, respectively. Alloys were then prepared using the same manufacturing process, and performance tests were conducted. The results confirmed that the alloys made from both coarse and fine powders possessed high tensile strength and excellent high-temperature creep properties. Their microstructure morphology, such as grain size, average size, distribution, and volume fraction of precipitates, was comparable. Furthermore, the average tensile strength and yield strength of the coarse and fine powder alloys differed by less than 10%, indicating that their overall mechanical properties were at the same level. This combination of powder particle sizes can meet different application requirements and significantly improve powder utilization. Previously, due to particle size selection, only 40-50% or even lower alloy powders could be utilized. Now, powders with particle sizes of 18-53 μm, commonly used in existing technologies, and powders with particle sizes of 53-150 μm discovered in this application, can be used to prepare target workpieces, increasing powder utilization to over 80% and significantly reducing costs.

[0028] Inclusions in powder primarily affect the fatigue properties of alloys, especially low-cycle fatigue properties. The workpieces prepared by the method provided in this application exhibit grain structure and multiple performance parameters comparable to existing nickel-based superalloys made from fine powder. They are particularly suitable for applications in hot-end components of aerospace, shipbuilding, and energy power systems, and possess strong tensile and creep resistance. This makes them especially suitable for components whose lifespan is limited by tensile or creep properties. The performance of these components is less constrained by fatigue performance but significantly limited by tensile and creep properties. The alloys prepared by the method of this application, when applied in these applications, not only meet performance requirements but also significantly improve the utilization rate of alloy powder. Attached Figure Description

[0029] To more clearly illustrate the technical solutions in the embodiments of this application or the prior art, the drawings used in the description of the embodiments or the prior art will be briefly introduced below. Obviously, the drawings described below are only some embodiments recorded in this application. For those skilled in the art, other drawings can be obtained based on these drawings without creative effort.

[0030] Figure 1 This is a comparison of the room temperature and high temperature tensile strength of the alloys in Examples 1-1 and 1-2 of the present invention; Figure 2 This is a comparison of the room temperature and high temperature yield strength of the alloys in Examples 1-1 and 1-2 of this invention; Figure 3 The tensile strength of the alloys in Example 1-1 and Comparative Examples 1 and 6 of this invention is compared under room temperature and high temperature conditions; wherein, the actual tensile temperature of the alloy in Example 1-1 is 760 ℃. Figure 4 The comparison shows the yield strength of the alloys in Example 1-1 and Comparative Examples 1 and 6 of this invention under room temperature and high temperature conditions; wherein, the actual tensile temperature of the alloy in Example 1-1 is 760 ℃. Figure 5 This is a comparison of the room temperature and high temperature tensile strength of the alloys in Examples 2-1 and 2-2 of the present invention; Figure 6 This is a comparison of the room temperature and high temperature yield strength of the alloys in Examples 2-1 and 2-2 of this invention; Figure 7 The tensile strength of the alloys in Example 2-1 and Comparative Examples 1 and 6 of this invention is compared under room temperature and high temperature conditions; wherein, the actual tensile temperature of the alloy in Example 2-1 is 760 ℃. Figure 8 The comparison shows the yield strength of the alloys in Example 2-1 and Comparative Examples 1 and 6 of this invention under room temperature and high temperature conditions; wherein, the actual tensile temperature of the alloy in Example 2-1 is 760 ℃. Figure 9 This is a comparison of the room temperature and high temperature tensile strength of the alloys in Examples 3-1 and 3-2 of this invention; Figure 10 This is a comparison of the room temperature and high temperature yield strength of the alloys in Examples 3-1 and 3-2 of this invention; Figure 11 The tensile strength of the alloys in Example 3-1 and Comparative Examples 1 and 6 of this invention is compared under room temperature and high temperature conditions; wherein, the actual tensile temperature of the alloy in Example 3-1 is 760 °C. Figure 12 The comparison shows the yield strength of the alloys in Example 3-1 and Comparative Examples 1 and 6 of this invention under room temperature and high temperature conditions; wherein, the actual tensile temperature of the alloy in Example 3-1 is 760 ℃. Figure 13 This is a comparison of creep strain time between Embodiments 1-1 and 1-2 of the present invention; Figure 14 This is a comparison of the creep strain time between Embodiment 2-1 and Embodiment 2-2 of the present invention; Figure 15 This is a comparison of the creep strain time between Embodiment 3-1 and Embodiment 3-2 of the present invention; Figure 16 This is a comparison of the grain structure of Examples 1-1 and 1-2 of the present invention after hot extrusion, where a is Example 1-1 and b is Example 1-2; Figure 17 The image shows a comparison of the grain structure of Examples 2-1 and 2-2 after hot extrusion, where a is Example 2-1 and b is Example 2-2. Figure 18 The image shows a comparison of the grain structure of Examples 3-1 and 3-2 after hot extrusion, where a is Example 3-1 and b is Example 3-2. Figure 19 The size and morphology of the primary γ' precipitate, secondary γ' precipitate, and carbide after hot extrusion in Examples 1-1 and 1-2 of the present invention are compared, where a is Example 1-1 and b is Example 1-2. Figure 20 The size and morphology of the primary γ' precipitate, secondary γ' precipitate, and carbide after hot extrusion in Examples 2-1 and 2-2 of the present invention are compared, where a is Example 2-1 and b is Example 2-2. Figure 21 The size and morphology of the primary γ' precipitate, secondary γ' precipitate, and carbide after hot extrusion in Examples 3-1 and 3-2 of the present invention are compared, where a is Example 3-1 and b is Example 3-2; Figure 22The image shows a comparison of the grain structure of Examples 1-1 and 1-2 of the present invention after extrusion and solution heat treatment, where a is Example 1-1 and b is Example 1-2. Figure 23 The image shows a comparison of the grain structure of Examples 2-1 and 2-2 of the present invention after extrusion and solution heat treatment, where a is Example 2-1 and b is Example 2-2. Figure 24 The image shows a comparison of the grain structure of Examples 3-1 and 3-2 of the present invention after extrusion and solution heat treatment, where a is Example 3-1 and b is Example 3-2; Figure 25 The image shows a comparison of the size and morphology of the secondary γ' precipitates and carbides after extrusion and solution heat treatment in Examples 1-1 and 1-2 of the present invention, where a is Example 1-1 and b is Example 1-2. Figure 26 The image shows a comparison of the size and morphology of the secondary γ' precipitates and carbides after extrusion and solution heat treatment in Examples 2-1 and 2-2 of the present invention, where a is Example 2-1 and b is Example 2-2. Figure 27 The image shows a comparison of the size and morphology of the secondary γ' precipitates and carbides after extrusion and solution heat treatment in Examples 3-1 and 3-2 of the present invention, where a is Example 3-1 and b is Example 3-2. Detailed Implementation

[0031] To enable those skilled in the art to better understand the technical solutions in this application, the technical solutions in the embodiments of this application will be clearly and completely described below. Obviously, the described embodiments are only a part of the embodiments of this application, and not all of the embodiments. Based on the embodiments in this application, all other embodiments obtained by those skilled in the art without creative effort are within the scope of protection of this application.

[0032] Representative formulations from this application were selected for experiments and compared with commonly used alloy FGH4097 formulations in the field and multiple formulations from patent CN110205523B related to this solution. The specific alloy formulations are shown in Table 1. Table 1 Alloy Composition

[0033] In Table 1, the alloy compositions of Examples 1-1 and 1-2 are the same, only the powder particle size is different, which is obtained by sieving the powder prepared by the atomization process. Examples 2-1 and 2-2, and Examples 3-1 and 3-2 are the same.

[0034] Alloy preparation The experimental alloys were prepared using the formulations of the examples and comparative examples in Table 1, according to the following method: S1. Prepare the master alloy by vacuum induction melting process using the above components, with a casting temperature of 1450 ℃. S2. The master alloy is made into powder by argon atomization process, and then the powder with the target particle size is obtained by sieving. S3. Process the above powders according to the same process: Pretreatment: cladding preparation, cladding welding, powder loading, degassing, and sealing welding. Then, the powder is densified using a hot isostatic pressing process. The hot isostatic pressing temperature is 1100℃, the pressure is 150MPa, and the time is 4 hours. The hot isostatically pressed alloy is machined to remove the outer cladding of the alloy. The straight cladding is completely removed. Then, a stainless steel cladding is used to wrap the hot isostatically pressed alloy ingot. Then, hot extrusion is performed to refine the grain structure. The hot extrusion temperature is 1080℃ and the extrusion ratio is 6:1. The hot-extruded alloy was subjected to solution heat treatment and aging heat treatment successively; among them... Solution heat treatment process: 1190℃, hold for 2 hours, then rapidly cool to room temperature; Aging heat treatment process: 760 ℃, hold for 8 hours, then rapidly air cool to room temperature.

[0035] The properties of the prepared samples were tested under the same conditions, and the results are as follows: I. Tensile Properties Table 2 compares the tensile strength and yield strength of the alloys in the examples and comparative examples.

[0036]

[0037] From Table 2, and the appendix Figures 1 to 12 It can be seen that the tensile properties of the coarse powder examples (Examples 1-1, 2-1 and 3-1) and the fine powder examples (Examples 1-2, 2-2 and 3-2) at different temperatures are at the same level, and the average tensile strength and yield strength differ by less than 10% of the average value, and are better than the alloy of Comparative Example 1.

[0038] II. Creep Properties The creep performance of each embodiment was tested, and the results are as follows: Table 3. Time to reach 0.2% creep strain for alloys with different powder particle sizes under different creep conditions

[0039] From Table 3, and the appendix Figures 13 to 15It can be seen that the creep properties of the coarse powder and fine powder alloys at different temperatures are excellent and at the same level.

[0040] III. Comparison of Grain Structure 3.1 The alloys prepared from coarse and fine powders in each embodiment were subjected to hot extrusion treatment, and the extruded grain structure was compared, such as... Figures 16 to 18 As shown.

[0041] in, Figure 16 In Example a, the average grain size after extrusion of coarse powder is 2.3 μm (particle size 53-150 μm); in Example b, the average grain size after extrusion of coarse powder is 2.6 μm (particle size 18-53 μm).

[0042] Figure 17 In Example 2-1 (particle size 53-150 μm), the average grain size after coarse powder extrusion is 1.2 μm; in Example 2-2 (particle size 18-53 μm), the average grain size after coarse powder extrusion is 1.1 μm.

[0043] Figure 18 In Example 3-1 (particle size 53-150 μm), the average grain size after extrusion of coarse powder is 1.3 μm; in Example 3-2 (particle size 18-53 μm), the average grain size after extrusion of coarse powder is 1.3 μm.

[0044] 3.2. The alloys prepared from coarse and fine powders in each embodiment were subjected to hot extrusion treatment. The microstructure of the extruded precipitates was compared, and the size and morphology of the primary γ' precipitates, secondary γ' precipitates, and carbides after extrusion were compared. Figures 19 to 21 As shown, although the particle size of the coarse powder example is different from that of the fine powder example, the grain size is almost the same after extrusion, and the size and volume fraction of the precipitated phase are geometrically identical.

[0045] 3.3 Furthermore, the grain structure of the alloys prepared from coarse and fine powders in each embodiment after solution heat treatment was compared, such as... Figures 22 to 24 As shown; in, Figure 22 In Example a, the average grain size after extrusion of coarse powder is 36.6 μm (particle size 53-150 μm); in Example b, the average grain size after extrusion of coarse powder is 29.9 μm (particle size 18-53 μm).

[0046] Figure 23 In Example 2-1 (particle size 53-150 μm), the average grain size after coarse powder extrusion is 20.6 μm; in Example 2-2 (particle size 18-53 μm), the average grain size after coarse powder extrusion is 18.6 μm.

[0047] Figure 24 In Example 3-1 (particle size 53-150 μm), the average grain size after coarse powder extrusion is 32.6 μm; in Example 3-2 (particle size 18-53 μm), the average grain size after coarse powder extrusion is 23.4 μm.

[0048] 3.4 Furthermore, a secondary comparison was made of the size and morphology of the γ' precipitate and carbides obtained after solution heat treatment in each embodiment, such as... Figures 25 to 27 As shown, comparing the coarse powder example and the fine powder example, although the powder particle size is different, after extrusion and solution heat treatment, the grain size difference is not significant (e.g., ASTM 5-9 grade), and the precipitated phase size and volume fraction are geometrically the same.

[0049] The above description of the disclosed embodiments enables those skilled in the art to make or use the invention. Various modifications to these embodiments will be readily apparent to those skilled in the art, and the general principles defined herein may be implemented in other embodiments without departing from the spirit or scope of the invention. Therefore, the invention is not to be limited to the embodiments shown herein, but is to be accorded the widest scope consistent with the principles and novel features disclosed herein.

Claims

1. A method of improving the utilization of a nickel-based powder superalloy, characterized by, The master alloy is subjected to a powdering process to obtain alloy powder, the alloy powder is screened to obtain powder with a particle size less than 150 μm, and the powder is divided into fine powder and coarse powder; the fine powder refers to powder with a particle size less than 53 μm, and the coarse powder refers to powder not less than 53 μm; the fine powder and the coarse powder are respectively subjected to a forming process to prepare a target alloy.

2. The method of claim 1, wherein, The powdering process is one of vacuum gas atomization, electrode induction gas atomization and rotating electrode atomization.

3. The gas atomizing powder production method according to claim 2, wherein The powdering process uses high-purity argon as atomization medium.

4. The method of claim 1, wherein, The alloy powder is screened to obtain fine powder with a particle size greater than or equal to 18 μm and less than 53 μm and coarse powder with a particle size greater than or equal to 53 μm and less than 150 μm, and the fine powder and the coarse powder are respectively subjected to a forming process to prepare a target alloy.

5. The method of claim 4, wherein, The coarse powder is subjected to a forming process to prepare a target alloy.

6. The method according to any one of claims 1-5, characterized in that, The method comprises the following steps: S1, preparing a master alloy by a vacuum induction melting process; S2, preparing alloy powder from the master alloy by a gas atomization process, and screening the alloy powder to obtain powder; S3, treating the powder by a hot isostatic pressing process; S4, performing hot deformation treatment on the alloy after hot isostatic pressing; S5, performing heat treatment on the alloy after hot deformation treatment.

7. The method of claim 6, wherein, Before the hot isostatic pressing in step S3, the method further comprises the steps of preparing a can, welding the can, loading powder, degassing and sealing the can; After step S3 and before step S4, the method further comprises the step of removing the can, and then performing external wrapping treatment on the hot isostatic pressed alloy ingot by using a stainless steel can; In step S4, the hot deformation treatment comprises any one or both of hot extrusion treatment and hot forging treatment; In step S5, the heat treatment comprises any one or both of solid solution heat treatment and aging heat treatment.

8. The method of claim 7, wherein, In step S1, the pouring temperature of the vacuum induction melting is 1400-1550 ℃; In step S3, the hot isostatic pressing temperature is 1050-1200 ℃, the pressure is 100-180 MPa, and the time is 3-8 hours; In step S4, the hot extrusion temperature is 1050-1150 ℃, and the extrusion ratio is 4-8; The hot forging temperature ranges from 1050 to 1150 ℃, and the rate is less than 0.01 / s; In step S5, the solid solution heat treatment is: heating to 1130-1200 ℃, holding for 1-4 hours, and then rapidly cooling to room temperature; The aging heat treatment is: heating to 750-800 ℃, holding for 6-12 hours, and then air cooling to room temperature.

9. The method of claim 1, wherein, The master alloy comprises the following components in mass fraction: Co 24%-30%; Cr 11%-14%; Mo 3.5%-5%; W 1%-2.8%; Al 2%-5%; Ti 2%-4%; Nb 1%-3%; Ta 2%-3.8%; C 0.02%-0.1%; B 0.01%-0.1%; Zr 0.02%-0.08%; the balance is Ni and impurities; or The master alloy comprises the following components in mass fraction: Co 10% to 14%; Cr 10% to 14%; Mo 3.5% to 5%; W 1% to 2.8%; Al 2% to 5%; Ti 2% to 4%; Nb 0.5% to 2.5%; Ta 4% to 7%; C 0.02% to 0.1%; B 0.02% to 0.1%; Zr 0.02% to 0.1%; balance Ni and impurities; or, The master alloy comprises the following components in mass fraction: Co 10% to 14%; Cr 10% to 14%; Mo 4% to 6%; W 0-1.5%; Al 2% to 4%; Ti 2% to 4%; Nb 0.75% to 1.4%; Ta 6.2% to 8.5%; C 0.11% to 0.2%; B 0.02% to 0.1%; Zr 0.02% to 0.1%; balance Ni and impurities.

10. The alloy prepared by the method of any one of claims 1 to 9 for use in preparing hot end components of aerospace, ship, energy power system.

Images