Nickel-based high-temperature alloy GH3536 and SLM preparation and heat treatment process thereof

Abstract

The invention relates to the field of high-temperature alloy material processing, in particular to a nickel-based high-temperature alloy GH3536 and an SLM preparation and heat treatment process thereof. Through the synergistic effect of ILR and HT, the core pain points of carbide precipitation imbalance and high-temperature instability coarsening in the GH3536 alloy formed through SLM are solved. A high-density subboundary created by the ILR drives short-distance diffusion of carbides at an HT stage to form a uniform spherical nano chain-shaped precipitated phase, so that the strength-plasticity inversion contradiction of a traditional process is broken through, and the strong-plasticity synergistic improvement of the tensile strength and the ductility is realized; the method comprises the steps that in-situ layer-by-layer laser remelting is conducted through the same technological parameters in the SLM process, and then the GH3536 high-temperature alloy is subjected to heat preservation in a vacuum furnace and cooled along with the furnace. The GH3536 high-temperature alloy prepared through the method has excellent performance, the plasticity of the GH3536 high-temperature alloy is equivalent to that of a forged piece, the strength of the GH3536 high-temperature alloy is higher than that of the forged piece, and the GH3536 high-temperature alloy has good industrial application

Description

Technical Field

[0001] The present invention relates to the field of high-temperature alloy material processing, and in particular to a nickel-based high-temperature alloy GH3536 and a SLM preparation and heat treatment process thereof. Background Art

[0002] As a typical nickel-based high-temperature alloy, GH3536 high-temperature alloy has excellent comprehensive properties, including high high-temperature strength, good oxidation resistance and hot corrosion resistance, which gives it broad application prospects in gas turbine blades, combustion chamber components and hot end components of aircraft engines.

[0003] It is well known that metals and alloys produced by selective laser melting (SLM) consist of large columnar grains with preferred orientation. This microstructure usually leads to undesirable anisotropic mechanical properties and is not conducive to the Hall-Page effect. When SLM is used to form GH3536 alloy, ultra-high-speed cooling (10 6 ~10 8 The high K / s (K / s) prevents elements such as Cr, Mo, and C from fully diffusing, resulting in a non-equilibrium supersaturated solid solution and loss of carbide strengthening. Furthermore, SLM-fabricated GH3536 alloys inevitably generate numerous microcracks, which negatively impacts material properties. Currently, GH3536 superalloys manufactured using selective laser melting (SLM) mostly follow the same heat treatment process as conventional manufacturing methods (such as casting and powder metallurgy). Due to fundamental differences in the manufacturing principles between SLM and conventional manufacturing methods, the microstructure of superalloys produced using this technology differs significantly from those produced using conventional methods. Furthermore, SLM-fabricated superalloys exhibit distinct mechanical properties compared to conventional methods. Conventional heat treatment requires prolonged, high-temperature treatment to eliminate columnar crystals, resulting in grain coarsening and strength loss. Hot isostatic pressing and post-processing methods are costly and uncontrollable in the morphology of the precipitated phases. Therefore, applying heat treatment processes suitable for conventional manufacturing processes to GH3536 superalloys manufactured using SLM may result in performance degradation, making them unsuitable for practical applications.

[0004] Therefore, it is urgent to develop a special heat treatment method that can target the unique microstructure and defect characteristics of SLM-formed GH3536 alloy, effectively refine the grains and eliminate microcracks, while promoting the full diffusion of elements such as Cr, Mo, and C and regulating the precipitation morphology of carbides, thereby achieving a comprehensive improvement in the mechanical properties of the material. Summary of the Invention

[0005] To achieve one of the above objectives, the present invention proposes a nickel-based high-temperature alloy GH3536 and its SLM preparation and heat treatment process. Through the "in-situ laser remelting (ILR) + vacuum insulation heat treatment (HT)" coupling process, the two-step synergistic effect solves the pain points of carbide precipitation imbalance and high-temperature instability. The method improves the performance of the GH3536 high-temperature alloy prepared by SLM and has good industrial application prospects. The technical solution of the present invention is achieved as follows:

[0006] In a first aspect, the present invention provides a SLM preparation and heat treatment process for a nickel-based high-temperature alloy GH3536, comprising the following steps:

[0007] S1. Dry and screen the initial powder of GH3536 high-temperature alloy, and start laser printing using an SLM printer under the protection of an inert gas; the inert gas includes argon with a purity greater than 99.99%;

[0008] S2. During the printing process, in-situ laser remelting is performed layer by layer. After printing is completed, the sample is separated from the substrate by wire cutting to obtain a preliminary sample. The substrate is made of 304 stainless steel, which is stress-relieved at 400-500°C before use and cleaned with alcohol to ensure no contamination.

[0009] S3. Soak the sample in alcohol and ultrasonically clean it three times, each time for 8 to 10 minutes, blow dry it, and then place it in a vacuum heating furnace for vacuum insulation heat treatment, and cool it in the furnace to obtain the GH3536 high-temperature alloy.

[0010] Preferably, the laser process parameters in step S2 are the same as those in step S1, and step S2 specifically includes: within one powder laying cycle, the laser scanning times are 2 times, and the scanning strategy is to rotate 67° layer by layer; rotating 67° layer by layer can disperse the direction of thermal stress and avoid stress concentration. Fixed scanning direction may lead to preferential growth of grains (such as <100> texture), while rotational scanning can weaken the texture effect and make the performance more uniform.

[0011] Further preferably, the second scan (remelting) can eliminate the porosity and unfused defects formed in the first scan, refine the grains, and improve the density. Only one remelting is enough to achieve a balance between efficiency and performance. If the number is increased to more than two times, it may lead to: excessive energy input, overheating of the molten pool, element burnout, grain coarsening and other problems, which in turn reduce the mechanical properties; multiple thermal cycles will intensify residual stress and increase the risk of part deformation or cracking.

[0012] Further preferably, the laser process parameters include: laser power of 220-280W, scanning speed of 600-800mm / s, laser scanning spacing of 80-100μm, powder layer thickness of 20-40μm, energy density of 107-291J / mm3 If the energy density is too low, incomplete fusion may occur easily; if the energy density is too high, coarse grains may occur easily, and the metal in the molten pool may vaporize and produce pores.

[0013] Preferably, in step S3, the holding temperature is 1150-1200° C., the holding time is 1-2 hours, and the heating rate is 5-8° C. / min.

[0014] Preferably, in step S3, the vacuum pressure is 10 -3 Below Pa.

[0015] Preferably, in step S1, the initial powder of the GH3536 high-temperature alloy includes the following components in mass percentage: Cr: 21.4%, Fe: 18.6%, Mo: 9%, Co: 1.5%, W: 0.58%, C: 0.09%, and the rest is Ni, and the sum of the mass percentages of the components is 100%.

[0016] Preferably, the initial powders are all spherical, with a particle size of 13 to 53 μm and a purity of >99.5%.

[0017] More preferably, in step S1, the initial powder needs to be vacuum dried at 60-80° C. for 8-10 hours.

[0018] Further preferably, in step S1, the initial powder needs to be sieved 2 to 3 times using a 250-mesh sieve.

[0019] In a second aspect, the present invention provides a nickel-based high-temperature alloy GH3536 obtained by the SLM preparation and heat treatment process described in the first aspect.

[0020] Compared with the prior art, the advantages of the present invention are:

[0021] (1) The present invention can improve work efficiency by in-situ layer-by-layer laser remelting + vacuum sintering during the SLM process, and reduce the cumbersome steps of traditional methods such as solution treatment followed by graded aging heat treatment, solution treatment + hot isostatic pressing post-treatment, etc.

[0022] (2) The present invention is directed to the SLM-formed GH3536 high-temperature alloy material, which significantly improves the plasticity of the material while ensuring the material strength, thereby broadening the industrial application prospects of SLM-formed GH3536. BRIEF DESCRIPTION OF THE DRAWINGS

[0023] In order to more clearly illustrate the embodiments of the present invention or the technical solutions in the prior art, the following briefly introduces the drawings required for use in the embodiments or the description of the prior art. Obviously, the drawings described below are only some embodiments of the present invention. For ordinary technicians in this field, other drawings can be obtained based on these drawings without paying any creative work.

[0024] Figure 1 This is a microstructure diagram of the GH3536 high-temperature alloy after remelting and heat treatment in Example 1 of the present invention;

[0025] Figure 2 (a) Precipitation phase morphology of the remelted and heat-treated GH3536 high-temperature alloy according to Example 1 of the present invention; (b) Elemental composition diagram;

[0026] Figure 3 This is a geometrically necessary dislocation density distribution diagram of the GH3536 high-temperature alloy after remelting in Example 1 of the present invention;

[0027] Figure 4 This is a geometrically necessary dislocation density distribution diagram of the remelted and heat-treated GH3536 high-temperature alloy according to Example 1 of the present invention;

[0028] Figure 5 This is a GOS map of the remelted and heat-treated GH3536 high-temperature alloy according to Example 1 of the present invention;

[0029] Figure 6 This is the metallographic structure diagram of the GH3536 high-temperature alloy prepared and processed in Comparative Example 1 of the present invention;

[0030] Figure 7 This is a morphology of the precipitation phase of the GH3536 high-temperature alloy prepared in Comparative Example 1 of the present invention;

[0031] Figure 8 This is the geometric necessary dislocation density distribution diagram of the GH3536 high-temperature alloy prepared and processed in Comparative Example 1 of the present invention. DETAILED DESCRIPTION

[0032] The embodiments of the present invention are described in detail below, and examples of the embodiments are shown in the accompanying drawings. The embodiments described below with reference to the accompanying drawings are exemplary and intended to be used to explain the present invention, but should not be understood as limiting the present invention.

[0033] It should be noted that the terms "first" and "second" are used for descriptive purposes only and should not be understood to indicate or imply relative importance or implicitly specify the number of the technical features indicated. Therefore, features defined as "first" or "second" may explicitly or implicitly include one or more of such features. Furthermore, in the description of the present invention, unless otherwise specified, "plurality" means two or more.

[0034] In this document, the terms “contain”, “include” or “include” are open expressions, that is, they include the contents specified in the present invention but do not exclude other contents.

[0035] As used herein, the terms "optional," "optionally," or "optionally" generally mean that the subsequently described event or circumstance may but need not occur, and that the description includes instances where the event or circumstance occurs and instances where it does not.

[0036] It should be noted that, although single in-situ laser remelting can refine grains and reduce porosity, it increases dislocation density, exacerbates residual stress accumulation, and carbides still exist in a microsegregated form, with most grains still columnar. Single heat treatment reduces dislocation density and residual stress, but it is difficult to interrupt the growth orientation of columnar crystals, resulting in uncontrolled carbide morphology, grain coarsening, and strength loss. The present invention's "in-situ laser remelting + vacuum insulation heat treatment" coupling process solves the pain points of carbide precipitation imbalance and high-temperature instability through a two-step synergistic effect.

[0037] The first step is in-situ layer-by-layer laser remelting: on the one hand, the solidified area is locally reheated through the second scan to make the molten pool flow again, filling the gaps and microcracks that were not completely fused in the first melting, thereby improving the density and uniformity, and promoting the carbide segregation elements Cr / Mo to re-dissolve in the matrix; on the other hand, it significantly refines the grains, reduces the dendrite spacing, and significantly increases the subgrain boundary and dislocation density, providing more nucleation sites, driving forces and short-range diffusion channels for subsequent element migration and recrystallization boundaries.

[0038] The second step is vacuum insulation heat treatment: it makes the subgrain boundaries and molten pool boundaries disappear, the grains are equiaxed, and the spherical carbide particles are promoted to precipitate in chains along the grain boundaries and nano-carbide particles are precipitated inside the grains. At the same time, it significantly eliminates small pores and dislocations, reduces the thermal stress of the material, and improves the strength and plasticity of the material.

[0039] The following is a clear and complete description of the technical solutions in the embodiments of the present invention, in conjunction with the accompanying drawings. Obviously, the described embodiments are only part of the embodiments of the present invention, not all of the embodiments. All other embodiments obtained by ordinary technicians in this field based on the embodiments of the present invention without making any creative efforts shall fall within the scope of protection of the present invention.

[0040] Example 1

[0041] This embodiment provides a SLM preparation and heat treatment process for a nickel-based high-temperature alloy GH3536, comprising the following steps:

[0042] S1. Commercially available GH3536 superalloy powder (spherical; particle size 13-53 μm; Cr: 21.4%, Fe: 18.6%, Mo: 9%, Co: 1.5%, W: 0.58%, C: 0.09%, and the remainder Ni) from Ningbo Zhongyuan New Material Technology Co., Ltd. was used as the initial powder. The powder was vacuum dried at 70°C for 8 h and sieved twice with a 250-mesh sieve.

[0043] S2. Add the mixed powder obtained in S1 into the material chamber of the SLM printing equipment, debug the parameters of the working chamber, and complete the preparation work before printing; under the protection of argon gas with a purity greater than 99.99%, use the SLM printer to start the laser printing. The specific process parameters are: laser power of 240W, scanning speed of 600mm / s, laser scanning spacing of 80μm, powder layer thickness of 30μm, energy density of 200J / mm 3 ;

[0044] S3. During the above printing process, in-situ laser remelting is required layer by layer, that is, remelting is performed with the same process parameters after each layer is printed. After printing is completed, the sample is separated from the substrate by wire cutting to obtain a preliminary sample; the substrate is stress-relieved at 500°C and cleaned with alcohol before use;

[0045] S4. Immerse the sample in alcohol, perform ultrasonic cleaning, and then blow dry. Then, place the sample in a vacuum heating furnace, heat it to 1175° C. at 5° C. / min, keep the temperature for 1 hour, and then cool it in the furnace to obtain the GH3536 high-temperature alloy.

[0046] After the alloy was processed in Example 1, the tensile strength of the material reached 960.2 MPa and the elongation was 46.6% when stretched at room temperature according to the national standard GB / T 228.1-2021. The strength exceeded the level of GH3536 forged by the traditional process, and the plasticity was equivalent to that of forged GH3536. The data of forged GH3536 was obtained from the literature (Mechanical property study on rapid additive layer manufacture). X alloy by selective laser melting technology), the alloy produced has a tensile strength of 767 MPa and an elongation of 44%.

[0047] At the same time, the alloy treated in Example 1 was subjected to microscopic characterization, such as Figure 1 As shown, it can be seen that the micro cracks of the material are completely eliminated, the density reaches 99.9%, and the columnar crystals are completely transformed into equiaxed crystals; Figure 2 As can be seen from (a) in the figure, the carbides precipitated at the grain boundaries in the material are completely spherical, and a large number of nano-carbide particles are also precipitated inside the grains. They are distributed along the grain boundaries in the form of short chain spheres, and the size distribution is uniform. The smooth edges of the carbides precipitated at the rounded grain boundaries can reduce the risk of stress concentration and are more conducive to maintaining plasticity than sharp strip-shaped carbides.

[0048] Figure 3 This is the geometrically necessary dislocation density distribution diagram of the material after only the first step of remelting treatment in Example 1. It can be seen that the dislocation density in the material has increased significantly and the number of substructures has increased significantly; and the high-density area of ​​geometrically necessary dislocations is the preferential position for recrystallization nucleation, which can reduce the recrystallization temperature or promote grain homogenization. Therefore, the structure after remelting provides sufficient nucleation cores, short-range diffusion channels and driving force for subsequent vacuum insulation recrystallization.

[0049] Figure 4 : This is the geometric necessary dislocation density distribution diagram of the alloy after the final holding heat treatment in Example 1. It can be seen that the geometric dislocation density has decreased significantly, indicating that the residual thermal stress in the material has been significantly eliminated; Figure 5 This is the GOS map of the remelted and heat-treated GH3536 superalloy, indicating more complete recrystallization. The GOS value reflects the degree of orientation variation within the grains and can be used to characterize the degree and uniformity of recrystallization. Lower GOS values ​​(blue and green) indicate high orientation consistency within the grains, typically corresponding to fully recrystallized grains. Higher GOS values ​​(red, yellow, and orange) indicate a large orientation gradient within the grains, indicating incomplete recrystallization and the potential for more dislocation entanglement and strain energy storage.

[0050] Example 2

[0051] This embodiment provides a SLM preparation and heat treatment process for a nickel-based high-temperature alloy GH3536, comprising the following steps:

[0052] S1. Commercially available GH3536 superalloy powder (spherical; particle size 13-53 μm; Cr: 21.4%, Fe: 18.6%, Mo: 9%, Co: 1.5%, W: 0.58%, C: 0.09%, and the remainder Ni) from Ningbo Zhongyuan New Material Technology Co., Ltd. was used as the initial powder. The powder was vacuum dried at 80°C for 8 h and sieved three times with a 250-mesh sieve.

[0053] S2. Add the mixed powder obtained in S1 into the material chamber of the SLM printing equipment, debug the parameters of the working chamber, and complete the preparation work before printing; under the protection of argon gas with a purity greater than 99.99%, use the SLM printer to start the laser printing. The specific process parameters are: laser power of 280W, scanning speed of 800mm / s, laser scanning spacing of 80μm, powder layer thickness of 40μm, energy density of 291J / mm 3 ;

[0054] S3. During the above printing process, in-situ laser remelting is required layer by layer, that is, remelting is performed with the same process parameters after each layer is printed. After printing is completed, the sample is separated from the substrate by wire cutting to obtain a preliminary sample; the substrate is stress-relieved at 500°C and cleaned with alcohol before use;

[0055] S4. Immerse the sample in alcohol, perform ultrasonic cleaning, and then blow dry. Then, place the sample in a vacuum heating furnace, heat it to 1175° C. at 5° C. / min, keep the temperature for 1.5 hours, and then cool it in the furnace to obtain the GH3536 high-temperature alloy.

[0056] After the alloy was processed according to the process of Example 2, the tensile strength of the material reached 960.2 MPa and the elongation was 46.6%.

[0057] Example 3

[0058] This embodiment provides a SLM preparation and heat treatment process for a nickel-based high-temperature alloy GH3536, comprising the following steps:

[0059] S1. Commercially available GH3536 superalloy powder (spherical; particle size 13-53 μm; Cr: 21.4%, Fe: 18.6%, Mo: 9%, Co: 1.5%, W: 0.58%, C: 0.09%, and the remainder Ni) from Ningbo Zhongyuan New Material Technology Co., Ltd. was used as the initial powder. The powder was vacuum-dried at 60°C for 10 h and sieved three times with a 250-mesh sieve.

[0060] S2. Add the mixed powder obtained in S1 into the material chamber of the SLM printing equipment, debug the parameters of the working chamber, and complete the preparation work before printing; under the protection of argon gas with a purity greater than 99.99%, use the SLM printer to start the laser printing. The specific process parameters are: laser power of 220W, scanning speed of 800mm / s, laser scanning spacing of 100μm, powder layer thickness of 20μm, energy density of 107J / mm 3 ;

[0061] S3. During the above printing process, in-situ laser remelting is required layer by layer, that is, remelting is performed with the same process parameters after each layer is printed. After printing is completed, the sample is separated from the substrate by wire cutting to obtain a preliminary sample; the substrate is stress-relieved at 600°C and cleaned with alcohol before use;

[0062] S4. Immerse the sample in alcohol, perform ultrasonic cleaning, and then blow dry. Then, place the sample in a vacuum heating furnace, heat it to 1200° C. at a rate of 5° C. / min, keep the temperature for 1.5 hours, and then cool it in the furnace to obtain the GH3536 high-temperature alloy.

[0063] After the alloy was processed according to the process of Example 3, the tensile strength of the material reached 873.2 MPa and the elongation was 48.3%.

[0064] Example 4

[0065] This embodiment provides a SLM preparation and heat treatment process for a nickel-based high-temperature alloy GH3536, comprising the following steps:

[0066] S1. Commercially available GH3536 superalloy powder (spherical; particle size 13-53 μm; Cr: 21.4%, Fe: 18.6%, Mo: 9%, Co: 1.5%, W: 0.58%, C: 0.09%, and the remainder Ni) from Ningbo Zhongyuan New Material Technology Co., Ltd. was used as the initial powder. The powder was vacuum dried at 70°C for 8 h and sieved twice with a 250-mesh sieve.

[0067] S2. Add the mixed powder obtained in S1 into the material chamber of the SLM printing equipment, debug the parameters of the working chamber, and complete the preparation work before printing; under the protection of argon gas with a purity greater than 99.99%, use the SLM printer to start the laser printing. The specific process parameters are: laser power of 240W, scanning speed of 600mm / s, laser scanning spacing of 80μm, powder layer thickness of 30μm, energy density of 200J / mm 3 ;

[0068] S3. During the above printing process, in-situ laser remelting is required layer by layer, that is, remelting is performed with the same process parameters after each layer is printed. After printing is completed, the sample is separated from the substrate by wire cutting to obtain a preliminary sample; the substrate is stress-relieved at 500°C and cleaned with alcohol before use;

[0069] S4. Immerse the sample in alcohol, perform ultrasonic cleaning, and then blow dry. Then, place the sample in a vacuum heating furnace, heat it to 1150° C. at 8° C. / min, keep the temperature for 2 h, and then cool it in the furnace to obtain the GH3536 high-temperature alloy.

[0070] After the alloy was processed according to the process of Example 4, the tensile strength of the material reached 873.2 MPa and the elongation was 48.3%.

[0071] Comparative Example 1

[0072] This comparative example provides an SLM preparation process for a nickel-based high-temperature alloy GH3536. The difference from Example 1 is that laser remelting is not performed in step S3, and the remaining steps are the same.

[0073] After the alloy was processed according to the process of Comparative Example 1, the tensile strength of the material reached 752.2 MPa and the elongation was 32.6%.

[0074] The sample formed without remelting process in Example 1 was found to have poor surface roughness and poor mechanical properties. Figure 6 It can be seen that there are still a large number of columnar crystals with consistent orientation in the material, which is anisotropic and is not conducive to the tensile properties of the material; Figure 7 It can be seen that the precipitated phase is irregular in shape and inconsistent in size, and there are relatively sharp carbides. The sharp edges of the irregular carbides will still become stress concentration points, especially under tensile loads, which may preferentially induce microcracks and then rapidly expand along the grain boundaries, leading to a tendency to brittle fracture; the carbides are irregular and long, which induce stress concentration, and there are many residual micropores and microcracks.

[0075] The precipitates in Example 1 are mainly spherical particles, and they show irregular semi-continuous shapes along the grain boundaries. The spherical carbides can relieve the stress in the grain boundary area, hinder the crack propagation, and also cause the crack to deflect and play a toughening role. At the same time, Figure 8 It can be seen that there are still a large number of geometrically necessary dislocations in the material of Comparative Example 1. On the one hand, this reflects that this treatment is not conducive to the recrystallization of the material. On the other hand, it may introduce long-range stress, which in turn causes the material to warp, crack or reduce fatigue life. The high dislocation density leads to severe work hardening, which is not conducive to plasticity.

[0076] Comparative Example 2

[0077] This comparative example provides an SLM preparation process for a nickel-based high-temperature alloy GH3536. The difference from Example 1 is that step S4 is omitted and vacuum insulation is not performed. The remaining steps are the same.

[0078] After the alloy was processed according to the process of Comparative Example 2, the tensile strength of the material reached 822.1 MPa and the elongation was 24.8%.

[0079] Similarly, the sample formed in Comparative Example 2 exhibited high residual stress, low strength, and poor plasticity. This is because the residual stress accumulation from remelting is more significant, and defects such as microcracks have not been completely eliminated, affecting the alloy's performance. Internal stress concentration and structural defects are detrimental to the material's plastic expansion.

[0080] Comparative Example 3

[0081] This comparative example provides a SLM preparation process for a nickel-based high-temperature alloy GH3536. The difference from Example 1 is that the vacuum insulation temperature is 1100° C., and the remaining steps are the same.

[0082] The alloy treated in Comparative Example 3 achieved a tensile strength of 798.2 MPa and an elongation of 42.3%. However, due to the presence of undissolved carbide strips (Cr / W enrichment), the recrystallization was insufficient, resulting in a large number of residual pores and insufficient improvement in strength and plasticity.

[0083] Comparative Example 4

[0084] This comparative example provides a SLM preparation process for a nickel-based high-temperature alloy GH3536. The difference from Example 1 is that the vacuum insulation temperature is 1220° C., and the remaining steps are the same.

[0085] The alloy treated by the process of Comparative Example 4 had a tensile strength of 801.7 MPa and an elongation of 38.5%. However, due to grain coarsening and carbide aggregation and growth, the strength and plasticity were not improved enough.

[0086] Comparative Example 5

[0087] This comparative example provides a SLM preparation process for a nickel-based high-temperature alloy GH3536. The difference from Example 1 is that the vacuum insulation time is 2.5 h, and the remaining steps are the same.

[0088] The alloy treated in Comparative Example 5 had a tensile strength of 735.6 MPa and an elongation of 35.2%. However, due to abnormal grain growth, carbides coarsened and continuously distributed along grain boundaries, resulting in insufficient improvement in strength and plasticity.

[0089] Comparative Example 6

[0090] This comparative example provides a SLM preparation process for a nickel-based high-temperature alloy GH3536. The difference from Example 1 is that the vacuum insulation time is 0.5 h, and the remaining steps are the same.

[0091] The alloy treated in Comparative Example 5 had a tensile strength of 724.7 MPa and an elongation of 36.4%. However, the alloy still had a large number of columnar grains, few and irregular precipitation surfaces, a large number of residual pores, and low density, resulting in insufficient improvement in strength and ductility.

[0092] The embodiments described above are some embodiments of the present invention, rather than all embodiments. The detailed description of the embodiments of the present invention is not intended to limit the scope of the invention as claimed, but merely represents selected embodiments of the present invention. All other embodiments obtained by those of ordinary skill in the art based on the embodiments of the present invention without creative work are within the scope of protection of the present invention.

Claims

1. A SLM preparation and heat treatment process for nickel-based high-temperature alloy GH3536, characterized in that: The following steps are involved: S1. Dry and screen the initial powder of GH3536 high-temperature alloy and perform laser printing; S2. During the above printing process, in-situ laser remelting is required layer by layer to obtain a preliminary sample; S3. Perform vacuum heat treatment on the sample to obtain the GH3536 high-temperature alloy.

2. The SLM preparation and heat treatment process of nickel-based high-temperature alloy GH3536 according to claim 1, characterized in that: The laser process parameters in step S2 are the same as those in step S1; step S2 specifically includes: within one powder spreading cycle, the laser scanning times is 2 times.

3. The SLM preparation and heat treatment process of the nickel-based high-temperature alloy GH3536 according to claim 2, characterized in that: The laser process parameters include: laser power of 220-280W, scanning speed of 600-800mm / s, laser scanning spacing of 80-100μm, powder layer thickness of 20-40μm, energy density of 107-291J / mm 3 .

4. The SLM preparation and heat treatment process of nickel-based high-temperature alloy GH3536 according to claim 1, characterized in that: In step S3, the holding temperature is 1150-1200° C., and the holding time is 1-2 hours.

5. The SLM preparation and heat treatment process of nickel-based high-temperature alloy GH3536 according to claim 1, characterized in that: In step S3, the vacuum pressure is 10 -3 Below Pa.

6. The SLM preparation and heat treatment process of nickel-based high-temperature alloy GH3536 according to claim 1, characterized in that: In step S1, the initial powder of the GH3536 high-temperature alloy includes the following components in mass percentage: Cr: 21.4%, Fe: 18.6%, Mo: 9%, Co: 1.5%, W: 0.58%, C: 0.09%, and the rest is Ni, and the sum of the mass percentages of the components is 100%.

7. The SLM preparation and heat treatment process of nickel-based high-temperature alloy GH3536 according to claim 1, characterized in that: In the step S1, the initial powders are all spherical, with a particle size of 13 to 53 μm and a purity of >99.5%.

8. The SLM preparation and heat treatment process of nickel-based high-temperature alloy GH3536 according to claim 1, characterized in that: In step S1, the initial powder needs to be vacuum dried at 60-80° C. for 8-10 hours.

9. The SLM preparation and heat treatment process of nickel-based high-temperature alloy GH3536 according to claim 1, characterized in that: In the step S1, the initial powder needs to be sieved 2 to 3 times using a 250-mesh sieve.

10. A nickel-based high-temperature alloy GH3536 obtained by the SLM preparation and heat treatment process according to claims 1 to 9.

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