A method for improving the intermediate temperature ductility of a wrought superalloy

Abstract

The application relates to a method for improving the medium-temperature plasticity of a wrought superalloy, which comprises the following steps: solid solution treatment is carried out on a formed wrought superalloy blank to obtain a solid solution treated blank; rough machining treatment is carried out on the solid solution treated blank to obtain a rough machining treated blank; the rough machining treated blank is annealed at a temperature of Tw+(10-30 DEG C) for 5-10 hours, and after cooling, a first aging treated blank is obtained; wherein Tw is the highest working temperature of the wrought superalloy part; the first aging treated blank is kept at a temperature of Tw+ / -10 DEG C for 25-100 hours, and after cooling, a second aging treated blank is obtained; and the second aging treated blank is subjected to finish machining treatment to obtain the wrought superalloy part. The application can effectively improve the grain boundary bonding force of the wrought superalloy, improve the plasticity of the alloy without changing the alloy composition, and provide a guarantee for the safe service of the wrought superalloy with high temperature bearing capacity.

Description

TECHNICAL FIELD

[0001] The present application relates to the technical field of wrought superalloys, and particularly relates to a method for improving the medium-temperature plasticity of wrought superalloys. BACKGROUND

[0002] Wrought superalloys are widely used in the technical fields of engines, power stations, nuclear energy and petrochemicals, etc. due to their excellent high-temperature strength, oxidation resistance, creep resistance and other comprehensive performances. The improvement of the performance of wrought superalloys is an important means to improve the thrust-to-weight ratio of engines, but in the design of wrought superalloys, researchers often focus on improving the high-temperature strength of the alloy and often ignore the thermal plasticity of the alloy. However, in the actual engineering preparation process, the tensile plasticity of the wrought superalloy will appear a plasticity trough phenomenon in the temperature range of 650-850 DEG C, and the tensile plasticity of the alloy will even be less than 5%. The temperature range just coincides with the service temperature range of the wrought alloy, so this plasticity trough phenomenon will damage the service performance of the wrought superalloy parts and is not conducive to the improvement of the performance of key components such as engines.

[0003] The medium-temperature brittleness is caused by the fact that the grain boundary strength is lower than the intragranular strength in a certain temperature range, so that the alloy is prone to intergranular cracking. The higher the temperature resistance of the wrought superalloy, the more precipitates contained in the intragranular, so the matching of the strength of the grain boundary and the intragranular is worse, and the alloy is more prone to intergranular cracking. Therefore, the main idea of improving the medium-temperature plasticity of the wrought superalloy with high temperature resistance is to improve the grain boundary strength of the alloy.

[0004] At present, the grain boundary is purified or the grain boundary binding force is increased by adding C, B, Zr or rare earth elements during alloy design, but if the trace elements are added too much, continuous carbides or borides will be easily formed at the grain boundary, which will deteriorate the performance of the alloy. During alloy smelting, the purity of the alloy melt is improved, and the contents of O, N and S inclusions in the alloy are reduced, but the wrought superalloy with high temperature resistance has adopted advanced triple smelting technology, and the content of inclusions in the alloy is controlled to be less than 30 ppm, and the inclusions in the alloy are already at a very low level.

[0005] In summary, there is an urgent need for an effective method to improve the grain boundary strength of the alloy to improve the medium-temperature plasticity of the wrought superalloy with high temperature resistance. SUMMARY

[0006] Therefore, the present application provides a method for improving the medium-temperature plasticity of wrought superalloys, and the main purpose is to make the wrought superalloy with high temperature resistance have excellent medium-temperature plasticity.

[0007] In order to achieve the above-mentioned purpose, the present application mainly provides the following technical scheme:

[0008] On one hand, embodiments of the present invention provide a method for improving the intermediate-temperature plasticity of deformed high-temperature alloys, comprising the following steps:

[0009] Solution treatment: The deformed high-temperature alloy billet after forming is subjected to solution treatment to obtain a solution-treated billet;

[0010] Rough machining: The billet after solution treatment is subjected to rough machining to obtain a rough-machined billet;

[0011] First aging treatment: The rough-machined billet is annealed at a temperature of Tw+(10~30℃) for 5~10h, and after cooling, the billet after the first aging treatment is obtained; wherein, Tw is the maximum working temperature of the deformed high-temperature alloy part;

[0012] Second aging treatment: The billet after the first aging treatment is kept at a temperature of Tw±10℃ for 25 to 100 hours, and after cooling, the billet after the second aging treatment is obtained.

[0013] Finishing process: The billet after the second aging treatment is finished to obtain deformed high-temperature alloy parts.

[0014] Preferably, the γ′ phase content in the deformed superalloy billet after the second aging treatment is less than 55 wt%.

[0015] Preferably, the γ′ phase content in the deformed superalloy billet after the second aging treatment is not less than 15 wt%.

[0016] Preferably, the γ′ phase size within the grains of the deformed high-temperature alloy billet after the second aging treatment is 50-200 nm.

[0017] Preferably, in the chemical composition of the deformed superalloy, the sum of the weight percentages of Al, Ti and Nb elements is ≥6wt%.

[0018] Preferably, in the chemical composition of the deformed superalloy, the sum of the weight percentages of W and Mo is ≥3.5wt%.

[0019] Preferably, in the chemical composition of the deformed high-temperature alloy, the weight percentage of Cr element is ≤20wt%.

[0020] Preferably, the deformed high-temperature alloy includes GH4065A alloy, GH4720Li alloy, and GH4068 alloy.

[0021] Preferably, in the solution treatment step: the shaped deformed superalloy billet is solution treated at a temperature of Tm (10-50°C) for 1-4 hours, and then cooled to room temperature to obtain the solution-treated billet; wherein, Tm is the complete dissolution temperature of the γ′ phase of the deformed superalloy; preferably, the cooling method is oil quenching or air cooling.

[0022] Preferably, the grain size of the billet after solution treatment is grade 7 or above.

[0023] Preferably, in the roughing process step: the solution-treated blank is machined to leave a margin of 1-5 mm.

[0024] Preferably, in the steps of the first aging treatment and the second aging treatment: the Tw is 700-800℃; and / or the cooling method is: air cooling or air quenching cooling.

[0025] Preferably, after the first aging treatment and the second aging treatment, the size of the γ′ phase in the deformed high-temperature alloy is 50-200 nm.

[0026] Preferably, after the second aging treatment, no TCP phase precipitation occurs in the microstructure of the deformed superalloy, and solid solution elements are segregated at the grain boundaries; discontinuous granular carbides and / or borides precipitate at the grain boundaries; preferably, the solid solution elements include W and Mo; preferably, the particle size of the granular carbides and / or borides is 0.5-3 μm.

[0027] Preferably, in the finishing process step: the blank after the second aging treatment is finished to achieve the target size, and the alloy part is shot peened to improve the surface properties.

[0028] On the other hand, embodiments of the present invention provide a deformed high-temperature alloy part, wherein no TCP phase precipitates in the microstructure of the deformed high-temperature alloy part, and solid solution elements are segregated at the grain boundaries; preferably, the solid solution elements include W and Mo; discontinuous granular carbides and / or borides precipitate at the grain boundaries of the deformed high-temperature alloy part.

[0029] Preferably, the deformed high-temperature alloy part is obtained by treating the formed deformed high-temperature alloy billet using any of the methods described above for improving the medium-temperature plasticity of the deformed high-temperature alloy.

[0030] Preferably, the deformed high-temperature alloy parts include GH4065A alloy parts, GH4720Li alloy parts, and GH4068 alloy parts.

[0031] Preferably, the tensile elongation of the deformed high-temperature alloy part at a temperature of 650–800°C is ≥10%, more preferably ≥20%, and even more preferably ≥28%.

[0032] Preferably, the tensile elongation of the deformed high-temperature alloy part at 700°C is ≥20%.

[0033] Preferably, the yield strength of the deformed high-temperature alloy part is ≥1000MPa at 700℃ and ≥900MPa at 800℃.

[0034] Preferably, the hardness of the deformed high-temperature alloy part is 35-50 HRC.

[0035] Compared with the prior art, the method of the present invention for improving the intermediate-temperature plasticity of deformed high-temperature alloys has at least the following beneficial effects:

[0036] The method for improving the intermediate-temperature plasticity of wrought superalloys provided in this invention includes the following steps: solution treatment of the formed wrought superalloy billet to obtain a solution-treated billet; rough machining of the solution-treated billet to obtain a rough-machined billet; annealing the rough-machined billet at a temperature of Tw + (10~30℃) for 5~10h, followed by cooling to obtain a billet after a first aging treatment; wherein Tw is the maximum working temperature of the wrought superalloy part; holding the billet after the first aging treatment at a temperature of Tw±10℃ for 25~100h, followed by cooling to obtain a billet after a second aging treatment; and finishing the billet after the second aging treatment to obtain a wrought superalloy part. It should be noted that the present invention effectively regulates elemental segregation and precipitate distribution at the alloy grain boundaries by controlling the temperature and time of heat treatment, thereby achieving the goal of improving the high-temperature grain boundary strength of the alloy. Specifically, (1) the purpose of solution treatment is to dissolve a large amount of the γ′ phase in the matrix, to make the alloy structure more uniform, to reduce the hardness of the alloy, to facilitate rough machining of the alloy, and to improve machining efficiency. The selection of solution temperature and holding time is based on the dissolution of the γ′ phase and the control of grain size, so a temperature of T is selected. mSolution treatment at (10-50℃) for 1-4 hours provides sufficient time for recrystallization and dissolution of the γ′ phase, while also inhibiting grain growth and preventing excessively large grain sizes. The hardness of the alloy after solution treatment is lower than HRC35. (2) The purpose of rough machining is to obtain near-net-shape dimensions, prevent the increase in hardness after aging treatment, increase the difficulty of machining, and improve tool life. This step should remove as much excess material as possible, and at the same time prevent deformation caused by excessively small billet size during subsequent aging treatment. Therefore, a 1-5 mm allowance is selected for the part. (3) The purpose of aging treatment is mainly to precipitate as much nano-sized γ′ phase as possible, so that the alloy has excellent high-temperature strength. This process should ensure that the alloy parts are structurally stable within the working temperature range, contain a high volume fraction of nano-sized γ′ phase, and strengthen the alloy. Therefore, it is first selected to be aging at T w Annealing at + (10~30℃) for 5-10h allows the γ′ phase-forming elements after solution treatment to precipitate, generating a high volume fraction of nanoscale γ′ phase within the grains. The purpose of step 4, aging treatment, is mainly to cause Mo and Wo atoms within the grains to segregate towards the grain boundaries, and even form granular carbides or borides with C or B atoms, thereby improving grain boundary bonding strength and enhancing the alloy's mid-temperature plasticity. Therefore, annealing at T... w The alloy was held at ±10℃ for 25–100 hours. After two aging processes, the size of the γ′ phase within the grains was between 50–200 nm, with no significant growth. Simultaneously, no TCP phase appeared in the alloy, and no continuous γ′ phase appeared at the grain boundaries. Step 5, finishing, aims to obtain the final shape and dimensions of the component. To improve the service life of the component, shot peening is performed on the finished component to create a compressive stress layer on its surface.

[0037] In summary, this invention utilizes heat treatment to induce the segregation of solid solution atoms at grain boundaries and to produce discontinuous granular carbides, thereby strengthening the grain boundaries and improving the mid-temperature plasticity of high-temperature deformable alloys with high temperature resistance. The strength of the alloy does not deteriorate significantly. The entire process involves few steps and is simple to operate. At the same time, it does not introduce any chemical reagents or toxic or polluting gases, making it environmentally friendly. Furthermore, this process is a combination of traditional processing methods. Most importantly, since this invention does not add any other alloying elements, it achieves the goal of optimizing alloy performance with almost no increase in alloy cost.

[0038] The above description is merely an overview of the technical solution of the present invention. In order to better understand the technical means of the present invention and to implement it in accordance with the contents of the specification, the preferred embodiments of the present invention are described in detail below with reference to the accompanying drawings. Attached Figure Description

[0039] Figure 1The diagram shows the change in element concentration distribution at the grain boundaries and the microstructure of the high-temperature-resistance deformable high-temperature alloy after treatment by the method of this invention.

[0040] Figure 2 The image shows the microstructure of the deformed high-temperature alloy after processing by the method described in Example 1 of this invention.

[0041] Figure 3 The image shows the microstructure of the deformed high-temperature alloy after being processed by the method described in Example 2 of this invention.

[0042] Figure 4 The image shows the microstructure of the deformed high-temperature alloy after being processed by the method described in Example 3 of this invention.

[0043] Figure 5 The image shows the microstructure of the deformed high-temperature alloy after being processed by the method described in Comparative Example 1.

[0044] Figure 6 The image shows the microstructure of the deformed superalloy after being processed by the method described in Comparative Example 2.

[0045] Figure 7 The image shows the microstructure of the deformed high-temperature alloy after being processed by the method described in Comparative Example 3. Detailed Implementation

[0046] To further illustrate the technical means and effects adopted by the present invention to achieve the intended purpose, the specific embodiments, structures, features, and effects according to the present invention will be described in detail below with reference to the accompanying drawings and preferred embodiments. In the following description, different "an embodiment" or "an embodiment" do not necessarily refer to the same embodiment. Furthermore, specific features, structures, or characteristics in one or more embodiments can be combined in any suitable form.

[0047] This invention utilizes a heat treatment method to cause solid solution strengthening elements in wrought superalloys to segregate towards grain boundaries, increasing grain boundary strength and further generating granular discontinuous carbides and / or borides, giving the alloy excellent mid-temperature plasticity. The method of this invention reduces or even eliminates waste caused by adding alloying elements and addresses the mid-temperature brittleness problem that traditional preparation methods cannot solve, thereby lowering the manufacturing cost of wrought superalloy parts and improving the safety factor of aero-engines.

[0048] The main solutions of this invention are as follows:

[0049] This invention provides a method for improving the intermediate-temperature plasticity of deformed high-temperature alloys, comprising the following steps:

[0050] Solution treatment: The deformed high-temperature alloy billet after forming is subjected to solution treatment to obtain a solution-treated billet.

[0051] It should be noted that "formed deformed high-temperature alloy billet" refers to the billet obtained by plastic deformation of deformed high-temperature alloy.

[0052] The specific steps are as follows: the shaped deformed superalloy billet is solution treated at a temperature of Tm (10-50℃) for 1-4 hours, and then cooled to room temperature to obtain the solution-treated billet; where Tm is the complete dissolution temperature of the γ′ phase of the deformed superalloy (Tm varies for each alloy and can generally be measured by DTA). The cooling method is oil quenching or air cooling.

[0053] The grain size of the billet after solution treatment is grade 7 or higher.

[0054] Rough machining: The billet after solution treatment is subjected to rough machining to obtain a rough-machined billet.

[0055] Specifically, this step involves machining the solution-treated blank, leaving a 1-5mm allowance for the machined part.

[0056] First aging treatment: The rough-machined billet is annealed at a temperature of Tw+(10~30℃) for 5~10h, and after cooling, the billet after the first aging treatment is obtained; wherein, Tw is the maximum working temperature of the deformed high-temperature alloy part.

[0057] The temperature Tw is 700–800℃. The cooling method is either air cooling or air quenching.

[0058] Second aging treatment: The billet after the first aging treatment is kept at a temperature of Tw±10℃ for 25 to 100 hours, and after cooling, the billet after the second aging treatment is obtained.

[0059] Here, Tw±10℃ refers to the temperature range of (Tw-10℃)~(Tw+10℃).

[0060] The temperature Tw is 700–800℃. The cooling method is either air cooling or air quenching.

[0061] Finishing process: The billet after the second aging treatment is finished to obtain deformed high-temperature alloy parts.

[0062] In this step, the blank after the second aging treatment is precision machined to achieve the target size, and the alloy part is shot peened to improve its surface properties.

[0063] The following should be noted regarding the above steps:

[0064] The method of this invention is aimed at high-temperature-resistance wrought high-temperature alloys, including but not limited to GH4065A, GH4720Li, and GH4068 alloys. First, the alloy billet is subjected to a T... m - Solution treatment at 10-50℃ for 1-4 hours is performed to maintain a uniform and fine alloy microstructure and reduce the alloy's hardness. After machining the billet into a near-target shape, it is then subjected to T... w Annealing at + (10~30℃) temperature for 5~10h allows uniform nanoscale γ′ phase to precipitate in the alloy, improving its high-temperature strength. Then, further annealing at T... w The alloy is held at ±10℃ for 25–100 hours to improve its grain boundary strength, giving it good plasticity in the mid-temperature range. Finally, it undergoes finishing to obtain the final shape and dimensions of the parts, and is shot-peened to obtain surface hardening and a compressive stress layer, thereby improving its service life.

[0065] Compared with the prior art, the present invention strengthens the grain boundaries by using a heat treatment method that is simple to operate, pollution-free and low-cost, without changing the degree of alloying of the material. This method involves the segregation of solid solution elements at the grain boundaries and the formation of discontinuous particulate carbides and borides, thereby improving the tensile plasticity of the alloy in the mid-temperature range.

[0066] In summary, this invention utilizes heat treatment to induce the segregation of solid solution atoms at grain boundaries and to produce discontinuous granular carbides, thereby strengthening the grain boundaries and improving the mid-temperature plasticity of high-temperature deformable alloys with high temperature resistance. The strength of the alloy does not deteriorate significantly. The entire process involves few steps and is simple to operate. At the same time, it does not introduce any chemical reagents or toxic or polluting gases, making it environmentally friendly. Furthermore, this process is a combination of traditional processing methods. Most importantly, since this invention does not add any other alloying elements, it achieves the goal of optimizing alloy performance with almost no increase in alloy cost.

[0067] The present invention will be further illustrated below with specific embodiments:

[0068] In embodiments of the present invention, the solution treatment temperature is between Tm and (10-50℃), the solution treatment time is 1-4 hours, followed by oil quenching or air cooling to room temperature; the solution-treated billet is then machined, retaining a 1-5 mm allowance; the rough-machined billet is then aged at Tw + (10-30℃) for 5-10 hours, followed by air cooling or gas quenching to room temperature; the billet is then aged at Tw ± 10℃ for 25-100 hours, followed by air cooling or gas quenching to room temperature; finally, the billet is finished to achieve the target dimensions, and the parts are shot-peened to improve surface properties. Furthermore, Auger electron spectroscopy is used to analyze the elemental concentration changes at grain boundaries, an Inspect F50 field emission scanning electron microscope is used to observe the alloy microstructure, and an INSTRON 5582 uniaxial tensile testing machine is used to test the alloy's plasticity and strength. The elemental concentration distribution at grain boundaries and microstructure photographs of the high-temperature-resistance deformable high-temperature alloy after treatment in the embodiments of the present invention are shown below. Figure 1 As shown, from Figure 1 It can be seen that solid solution elements segregate at the grain boundaries, and discontinuous carbides appear. The performance of the high-temperature resistant deformable high-temperature alloy treated according to the embodiments of the present invention is compared with that of the alloy treated in the comparative example, as detailed in Table 1.

[0069] Example 1

[0070] This embodiment improves the mid-temperature plasticity of GH4065A alloy (alloy blades) and mainly includes the following steps:

[0071] Solution treatment: The formed blade blank was solution treated at 1080°C for 1 hour, and then air-cooled to room temperature to obtain the solution-treated blank. Additionally, the Tm of the alloy in this embodiment was tested to be 1110°C.

[0072] Rough machining: The solution-treated blank is machined by turning, leaving a 1mm allowance for the part.

[0073] First aging treatment: The rough-machined billet is annealed at 760℃ for 10 hours, and then air-cooled to room temperature.

[0074] Second aging treatment: The billet after the first aging treatment is held at 750℃ for 50 hours and then air-cooled to room temperature. The γ′ phase size in the deformed superalloy billet after the second aging treatment is 50-200 nm.

[0075] Finishing process: The aging-treated blank is finished to achieve the target size, and the parts are shot-peened to improve surface performance.

[0076] The microstructure at the grain boundaries of the deformed superalloy part (blade) obtained in this embodiment is as follows: Figure 2 As shown. FromFigure 2 It can be seen that the average grain size of the alloy is 9 μm, and a large number of discontinuous carbides (particle size of 0.5-3 μm) are distributed on the grain boundaries of the alloy, thereby improving the grain boundary bonding strength and improving the medium-temperature plasticity of the alloy.

[0077] The properties of the deformed high-temperature alloy part (blade) prepared in this embodiment are shown in Table 1. Its yield strength at 700℃ is 1050MPa, tensile strength is 1250MPa, and elongation is 35%.

[0078] Comparative Example 1

[0079] Comparative Example 1 describes the preparation of a GH4065A alloy blade, which mainly includes the following steps:

[0080] Solution treatment: The formed blade blank is solution treated at 1080℃ for 1 hour and then air-cooled to room temperature.

[0081] Rough machining: The solution-treated blank is machined by turning, leaving a 1mm allowance for the part.

[0082] Aging treatment: The rough-machined billet is annealed at 760℃ for 10 hours, and then air-cooled to room temperature.

[0083] Finishing: The aging-treated billet is finished to achieve the target dimensions.

[0084] The microstructure at the grain boundaries of the deformed superalloy part prepared in Comparative Example 1 is as follows: Figure 5 As shown. From Figure 5 It can be seen that although the average grain size of the alloy is also around 9μm, no carbide precipitation occurs at the grain boundaries.

[0085] The properties of the deformed high-temperature alloy parts prepared in Comparative Example 1 are shown in Table 1. The yield strength at 700℃ is 1080 MPa, the tensile strength is 1297 MPa, and the elongation is 15%.

[0086] Example 2

[0087] This embodiment improves the mid-temperature plasticity of GH4720Li alloy turbine disks, mainly including the following steps:

[0088] Solution treatment: The formed turbine disk blank was solution treated at 1100°C for 4 hours, and then oil quenched to room temperature. Additionally, the Tm of the alloy in this embodiment was tested to be 1150°C.

[0089] Rough machining: The solution-treated blank is machined by turning, leaving a 5mm allowance for the part.

[0090] First aging treatment: The rough-machined billet is annealed at 760℃ for 8 hours, and then air-cooled to room temperature.

[0091] Second aging treatment: The billet after the first aging treatment is held at 750℃ for 50 hours and then air-cooled to room temperature. The γ′ phase size in the deformed superalloy billet after the second aging treatment is 50-200 nm.

[0092] Finishing process: The blank after the second aging treatment is finished to achieve the target size, and the parts are shot peened to improve surface performance.

[0093] The microstructure at the grain boundaries of the deformed high-temperature alloy part (turbine disk) obtained in this embodiment is as follows: Figure 3 As shown. From Figure 3 It can be seen that the average grain size of the alloy is 6 μm, and a large number of discontinuous carbides (particle size 0.5-3 μm) are distributed on the grain boundaries, thereby improving the grain boundary bonding strength and the mid-temperature plasticity of the alloy. The properties of the deformed high-temperature alloy part (turbine disk) obtained in this embodiment are shown in Table 1, with a yield strength of 1060 MPa at 700℃, a tensile strength of 1280 MPa, and an elongation of 30%.

[0094] Comparative Example 2

[0095] Comparative Example 2 involves the preparation of a GH4720Li alloy turbine disk, comprising the following steps:

[0096] Solution treatment: The formed turbine disk blank is solution treated at 1100℃ for 4 hours and then oil quenched to room temperature.

[0097] Rough machining: The solution-treated blank is machined by turning, leaving a 5mm allowance for the part.

[0098] First aging treatment: The rough-machined billet is annealed at 650℃ for 24 hours, and then air-cooled to room temperature.

[0099] Second aging treatment: The billet after the first aging treatment is kept at 760℃ for 16 hours and then air-cooled to room temperature.

[0100] Finishing process: The billet after the second aging treatment is finished to achieve the target size.

[0101] The microstructure at the grain boundaries of the deformed high-temperature alloy part (GH4720Li alloy turbine disk) obtained in Comparative Example 2 is as follows: Figure 6 As shown, from Figure 6It can be seen that although the average grain size of the alloy is also around 6 μm, no carbide precipitation occurs at the grain boundaries. The properties of the deformed high-temperature alloy part (GH4720Li alloy turbine disk) obtained in Comparative Example 2 are shown in Table 1. Its yield strength at 700℃ is 1070 MPa, its tensile strength is 1310 MPa, and its elongation is 18%.

[0102] Example 3

[0103] This invention provides a method for improving the mid-temperature plasticity of GH4068 alloy turbine disks, comprising the following steps:

[0104] Solution treatment: The formed turbine disk blank was solution treated at 1100℃ for 4 hours, and then oil quenched to room temperature. The Tm of the alloy in this embodiment was tested to be 1141℃.

[0105] Rough machining: The solution-treated blank is machined by turning, leaving a 5mm allowance for the part.

[0106] First aging treatment: The rough-machined billet is annealed at 760℃ for 10 hours, followed by air cooling to room temperature. The γ′ phase size within the grains of the deformed superalloy billet after the second aging treatment is 50-200 nm.

[0107] Second aging treatment: The billet after the first aging treatment is kept at 750℃ for 50 hours and then air-cooled to room temperature.

[0108] Finishing process: The blank after the second aging treatment is finished to achieve the target size, and the parts are shot peened to improve surface performance.

[0109] The microstructure at the grain boundaries of the deformed high-temperature alloy part (GH4068 alloy turbine disk) obtained in this embodiment is as follows: Figure 4 As shown, from Figure 4 It can be seen that the average grain size of the alloy is 5 μm, and a large number of discontinuous carbides (particle size 0.5-3 μm) are distributed on the grain boundaries, thereby improving the grain boundary bonding strength and the mid-temperature plasticity of the alloy. The properties of the deformed high-temperature alloy part (GH4068 alloy turbine disk) obtained in this embodiment are shown in Table 1. Its yield strength at 700℃ is 1100 MPa, its tensile strength is 1250 MPa, and its elongation is 28%.

[0110] Comparative Example 3

[0111] Comparative Example 3 describes the preparation of a GH4068 alloy turbine disk, including the following steps:

[0112] Solution treatment: The formed turbine disk blank is solution treated at 1100℃ for 4 hours and then oil quenched to room temperature.

[0113] Rough machining: The solution-treated blank is machined by turning, leaving a 5mm allowance for the part.

[0114] First aging treatment: The rough-machined billet is annealed at 650℃ for 24 hours, and then air-cooled to room temperature.

[0115] Second aging treatment: The billet after the first aging treatment is kept at 760℃ for 16 hours and then air-cooled to room temperature.

[0116] Finishing process: The billet after the second aging treatment is finished to achieve the target size.

[0117] The microstructure at the grain boundaries of the deformed high-temperature alloy part (GH4068 alloy turbine disk) obtained in Comparative Example 3 is as follows: Figure 7 As shown. From Figure 7 It can be seen that although the average grain size of the alloy is also around 5 μm, no carbide precipitation occurs at the grain boundaries. Figure 7 The white, bright substance in the sample is not a carbide, but merely a grain boundary (as confirmed by energy dispersive spectroscopy). The properties of the deformed high-temperature alloy part (GH4068 alloy turbine disk) obtained in Comparative Example 3 are shown in Table 1. Its yield strength at 700℃ is 1130 MPa, tensile strength is 1290 MPa, and elongation is 15%.

[0118] The properties of the deformed high-temperature alloy parts obtained after treatment in Examples 1-3 and Comparative Examples 1-3 are shown in Table 1.

[0119] Table 1

[0120]

[0121] As shown in Table 1, the processing method of this invention improves the mid-temperature plasticity of wrought superalloys (significantly increasing elongation) without changing the alloy composition, thus ensuring the safe service of high-temperature-resistant wrought superalloys. Furthermore, the method of this invention is simple to operate, low in cost, and environmentally friendly, which facilitates the application of high-temperature-resistant wrought superalloys in high-temperature structural components.

[0122] The above description is merely a preferred embodiment of the present invention and is not intended to limit the present invention in any way. Any simple modifications, equivalent changes, and alterations made to the above embodiments based on the technical essence of the present invention shall still fall within the scope of the technical solution of the present invention.

Claims

1. A method for improving the intermediate-temperature plasticity of deformed high-temperature alloys, characterized in that, It includes the following steps: Solution treatment: The deformed high-temperature alloy billet after forming is subjected to solution treatment to obtain a solution-treated billet; Rough machining: The billet after solution treatment is subjected to rough machining to obtain a rough-machined billet; First aging treatment: The rough-machined billet is annealed at a temperature of Tw+ (10~30℃) for 5~10h, and after cooling, the billet after the first aging treatment is obtained; wherein, Tw is the maximum working temperature of the deformed high-temperature alloy part; wherein, Tw is 700~800℃; Second aging treatment: The billet after the first aging treatment is kept at a temperature of Tw±10℃ for 25~100h, and after cooling, the billet after the second aging treatment is obtained. Finishing process: The billet after the second aging treatment is finished to obtain deformed high-temperature alloy parts; Among them, after the first aging treatment and the second aging treatment, the size of the γ′ phase in the deformed high temperature alloy is 50~200nm; The second aging treatment results in the absence of TCP phase precipitation in the microstructure of the deformed superalloy, with solid solution elements segregating at the grain boundaries; discontinuous granular carbides and / or borides precipitate at the grain boundaries; wherein the solid solution elements include W and Mo; and wherein the particle size of the granular carbides and / or borides is 0.5-3 μm. The deformed high-temperature alloys include GH4065A alloy, GH4720Li alloy, and GH4068 alloy.

2. The method for improving the intermediate-temperature plasticity of deformed high-temperature alloys according to claim 1, characterized in that, The γ′ phase content in the deformed superalloy billet after the second aging treatment is less than 55 wt%; and / or The γ′ phase content in the deformed superalloy billet after the second aging treatment is not less than 15wt%.

3. The method for improving the intermediate-temperature plasticity of deformed high-temperature alloys according to claim 1, characterized in that, In the chemical composition of the wrought superalloy: The sum of the weight percentages of Al, Ti, and Nb elements is ≥6 wt%; and / or The sum of the weight percentages of W and Mo is ≥3.5 wt%; and / or The weight percentage of Cr is ≤20wt%.

4. The method for improving the intermediate-temperature plasticity of deformed high-temperature alloys according to claim 1, characterized in that, In the solution treatment step: The shaped high-temperature alloy billet is subjected to solution treatment at a temperature of Tm (10~50℃) for 1~4h, and then cooled to room temperature to obtain the solution-treated billet; where Tm is the complete dissolution temperature of the γ′ phase of the shaped high-temperature alloy.

5. The method for improving the intermediate-temperature plasticity of deformed high-temperature alloys according to claim 4, characterized in that, The cooling method is either oil quenching or air cooling.

6. The method for improving the intermediate-temperature plasticity of deformed high-temperature alloys according to claim 1, characterized in that, The grain size of the billet after solution treatment is grade 7 or higher.

7. The method for improving the intermediate-temperature plasticity of deformed high-temperature alloys according to claim 1, characterized in that, In the roughing process step: The solution-treated blank is then machined, leaving a margin of 1-5 mm.

8. The method for improving the intermediate-temperature plasticity of deformed high-temperature alloys according to claim 1, characterized in that, In the steps of the first and second timeliness processing: The cooling method is either air cooling or air quenching cooling.

9. The method for improving the intermediate-temperature plasticity of deformed high-temperature alloys according to claim 1, characterized in that, In the finishing process step: The blank after the second aging treatment is precision machined to achieve the target size, and the alloy parts are shot peened to improve surface properties.

10. A deformable high-temperature alloy part, characterized in that, The microstructure of the deformed superalloy part has no TCP phase precipitation, but solid solution elements are segregated at the grain boundaries; wherein, the solid solution elements include W and Mo; discontinuous granular carbides and / or borides precipitate at the grain boundaries of the deformed superalloy part; wherein, the deformed superalloy part is obtained by treating the formed deformed superalloy billet with the method for improving the intermediate-temperature plasticity of the deformed superalloy as described in any one of claims 1-9.

11. The deformable high-temperature alloy part according to claim 10, characterized in that, The tensile elongation of the deformed high-temperature alloy part is ≥10% at a temperature of 650~800℃.

12. The deformable high-temperature alloy part according to claim 11, characterized in that, The tensile elongation of the deformed high-temperature alloy part is ≥20% at a temperature of 650~800℃.

13. The deformable high-temperature alloy part according to claim 12, characterized in that, The tensile elongation of the deformed high-temperature alloy part is ≥28% at a temperature of 650~800℃.

14. The deformable high-temperature alloy part according to claim 11, characterized in that, The tensile elongation of the deformed high-temperature alloy part at 700℃ is ≥20%.

15. The deformable high-temperature alloy part according to claim 11, characterized in that, The yield strength of the deformed high-temperature alloy part is ≥1000MPa at 700℃ and ≥900MPa at 800℃.

16. The deformed high-temperature alloy part according to claim 11, characterized in that, The hardness of the deformed high-temperature alloy part is 35~50HRC.

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