Nickel-based wrought superalloy, nickel-based wrought superalloy part and preparation method thereof

Abstract

The invention relates to a nickel-based wrought superalloy, a nickel-based wrought superalloy part and a preparation method of the nickel-based wrought superalloy part, and relates to the technical field of superalloy design and preparation. According to the main technical scheme, the nickel-based wrought superalloy comprises the following chemical components in percentage by weight: 20 to 30 percent of Co, 12 to 15 percent of Cr, 4.0 to 6.0 percent of Ti, 2.0 to 3.0 percent of Al, 0.5 to 3.0 percent of W, 2.0 to 4.0 percent of Mo, 0.01 to 0.08 percent of C, less than or equal to 2.0 percent of Nb, less than or equal to 2.0 percent of Ta, less than or equal to 0.06 percent of Zr, less than or equal to 0.05 percent of B and the balance of Ni. According to the preparation method, a mother alloy cast ingot meeting the requirements of the chemical components is prepared mainly through vacuum induction, electroslag remelting and vacuum self-consuming processes in a smelting mode, then diffusion annealing and forging treatment are carried out to prepare a bar with the uniform structure, the bar is cut out and formed, finally, solid solution aging treatment is carried out, and the nickel-based wrought superalloy part with the long service life is obtained. And the highest service temperature can reach 800 DEG C or above.

Description

Technical Field

[0001] This invention relates to the field of high-temperature alloy design and preparation technology, and in particular to a nickel-based wrought high-temperature alloy, nickel-based wrought high-temperature alloy parts, and their preparation methods. Background Technology

[0002] Wrought superalloys are key materials for the manufacture of power systems such as aero engines, gas turbines and generators. Existing aero-engine alloys are difficult to meet the comprehensive performance requirements of high temperature resistance, long service life and easy formability.

[0003] Traditional designs achieve high-temperature strength by increasing the degree of alloying. For example, the GH4151 alloy used in aero-engine turbine disks, with a temperature resistance of 800℃, contains elements such as Al, Ti, and Nb up to 10%, a γ′ phase volume fraction of approximately 55%, and a carbon content of approximately 0.1%. However, high alloying leads to coarsening of the γ′ phase and grain boundary carbide film formation during long-term service, causing instability in microstructure and properties, and making hot working difficult. Although the creep rupture life of long-life deformation superalloys used in ultra-supercritical applications can be extrapolated to 100,000 hours, the temperature resistance of these alloys is less than 700℃.

[0004] To improve the heat resistance of wrought superalloys, it is usually necessary to increase the content of precipitation strengthening, solid solution strengthening, and grain boundary strengthening elements. While high alloying can significantly improve high-temperature strength, it can also lead to excessive γ′ phase content, narrowing the matrix channels and shortening the element diffusion distance, thereby accelerating the coarsening and instability of the γ′ phase. Furthermore, when the temperature exceeds the isobaric temperature, the grain boundary strength is lower than that within the grains, easily inducing intergranular cracking. Therefore, increasing the content of grain boundary carbides is often used to strengthen the grain boundaries; however, this strategy can also cause the carbides to gradually thin during long-term service, thus deteriorating the alloy properties.

[0005] In summary, there is an urgent need for a wrought high-temperature alloy that not only has excellent strength and service life at 800~850℃, but also good high-temperature machinability, and can be used to manufacture core hot-end components in new engines and generators. Summary of the Invention

[0006] In view of this, the present invention provides a nickel-based wrought superalloy, nickel-based wrought superalloy parts and a method for preparing the same, the main purpose of which is to enable the nickel-based wrought superalloy to have excellent strength and service life at 800~850℃, while also having good high-temperature machinability.

[0007] To achieve the above objectives, the present invention mainly provides the following technical solutions:

[0008] On one hand, embodiments of the present invention provide a nickel-based wrought superalloy, wherein the chemical composition of the nickel-based wrought superalloy, by weight percentage, is as follows:

[0009] Co: 20~30wt%, Cr: 12~15wt%, Ti: 4.0~6.0wt%, Al: 2.0~3.0wt%, W: 0.5~3.0wt%, Mo: 2.0~4.0wt%, C: 0.01~0.08wt%, Nb≤2.0wt%, Ta≤2.0wt%, Zr≤0.06wt%, B≤0.05wt%, Ni is the balance.

[0010] Preferably, the chemical composition of the nickel-based wrought superalloy, by weight percentage, is as follows:

[0011] Co: 22.5~26.5wt%, Cr: 13.5~15.5wt%, Ti: 4.25~6.0wt%, Al: 2.2~2.6wt%, W: 1.1~1.5wt%, Mo: 2.5~3.5wt%, C: 0.01~0.04wt%, Nb: 0.5~1.0wt%, Ta: 0.5~1.0wt%, Zr: 0.025~0.04wt%, B: 0.01~0.04wt%, Ni is the balance.

[0012] Preferably, in the nickel-based wrought superalloy: the sum of the weight percentages of Al, Ti, Nb, and Ta is greater than or equal to 7.0 wt% and less than or equal to 9.5 wt%; and / or the sum of Nb and Ta is less than or equal to 2 wt%; and / or the sum of W and Mo is greater than or equal to 3% and less than or equal to 6 wt%.

[0013] On the other hand, embodiments of the present invention provide a nickel-based wrought high-temperature alloy part, wherein the material of the nickel-based wrought high-temperature alloy part is any of the nickel-based wrought high-temperature alloys described above; preferably, the nickel-based wrought high-temperature alloy part is a hot-end component.

[0014] Preferably, the γ´ phase content in the alloy microstructure of the nickel-based wrought superalloy part is 30-45 wt%, and the γ´ phase exhibits a bimodal distribution of 200-500 nm and <100 nm; and / or the morphology of the γ´ phase in the alloy microstructure of the nickel-based wrought superalloy part is spherical particles; and / or the average grain size of the alloy of the nickel-based wrought superalloy part is grade 4-6; and / or the yield strength of the nickel-based wrought superalloy part at 800℃ is ≥800MPa, the tensile strength is ≥900MPa, and the elongation is ≥5%; and / or the creep rupture life of the nickel-based wrought superalloy part under tensile stress conditions of 150MPa at 800℃ is ≥2000h; and / or the creep rupture life of the nickel-based wrought superalloy part under tensile stress conditions of 100MPa at 850℃ is ≥2000h.

[0015] Furthermore, embodiments of the present invention provide a method for preparing the above-mentioned nickel-based deformed high-temperature alloy part, which includes the following steps:

[0016] Step 1): The raw materials are subjected to vacuum induction melting, electroslag remelting, and vacuum consumable melting to obtain nickel-based deformed high-temperature alloy ingots;

[0017] Step 2): The nickel-based deformed superalloy ingot is subjected to diffusion annealing to obtain a diffusion-annealed ingot;

[0018] Step 3): The ingot after diffusion annealing is subjected to multi-fire forging to obtain nickel-based deformed high-temperature alloy bars;

[0019] Step 4): The nickel-based wrought high-temperature alloy bar is formed to obtain a shaped part;

[0020] Step 5): Perform solution treatment on the molded part to obtain the solution-treated molded part;

[0021] Step 6): The solution-treated formed part is subjected to aging treatment to obtain a nickel-based deformed high-temperature alloy part.

[0022] Preferably, in step 1): the raw material is subjected to vacuum induction melting to obtain a vacuum induction melt ingot; the vacuum induction melt ingot is first subjected to a first stress-relief annealing treatment and then subjected to electroslag remelting treatment; the electroslag remelted ingot obtained after electroslag remelting treatment is first subjected to a second stress-relief annealing treatment and then subjected to vacuum arc remelting treatment to obtain a nickel-based deformed high-temperature alloy ingot; preferably, the temperature of the first stress-relief annealing treatment is 900-1050℃, and the time of the first stress-relief annealing treatment is 5-15h; preferably, the temperature of the second stress-relief annealing treatment is 900-1050℃, and the time of the second stress-relief annealing treatment is 5-15h.

[0023] Preferably, in step 2), the nickel-based deformed high-temperature alloy ingot is heated to 1140-1170℃, held for 20-50 hours, then heated to 1170-1200℃, held for 20-50 hours, and cooled to obtain the ingot after diffusion annealing.

[0024] Preferably, in step 3), the diffusion-annealed ingot is subjected to multiple forging processes at a temperature of T - (10~40) ℃ to T + (10~30) ℃; wherein T is the complete dissolution temperature of the γ´ phase of the nickel-based wrought superalloy; preferably, the first 5~10 forging processes are carried out in the single-phase region to reduce deformation resistance and promote recrystallization, and the last 5~10 forging processes are carried out in the two-phase region; preferably, the temperature of the single-phase region is T to T + (10~30) ℃; and the temperature of the two-phase region is T - (10~40) ℃ to T.

[0025] Preferably, in step 4), the forming temperature is T - (10~40)℃, where T is the complete dissolution temperature of the γ' phase of the nickel-based wrought superalloy; and / or in step 5), the solution treatment temperature is (T-20)℃~(T+10)℃, and the solution treatment time is 2-6h, where T is the complete dissolution temperature of the γ' phase of the nickel-based wrought superalloy; and / or in step 6), the formed part after solution treatment is heated to 1000-1100℃ and held for 2-6h for a first aging treatment, and after cooling, a formed part after the first aging treatment is obtained; the formed part after the first aging treatment is heated to 800-860℃ and held for 10-24h for a second aging treatment, and after cooling, a nickel-based wrought superalloy part is obtained.

[0026] Compared with the prior art, the nickel-based wrought superalloy, the nickel-based wrought superalloy parts and the preparation method thereof of the present invention have at least the following beneficial effects:

[0027] On one hand, embodiments of the present invention provide a nickel-based wrought superalloy, wherein the chemical composition of the nickel-based wrought superalloy is as follows: Co: 20~30wt%, Cr: 12~15wt%, Ti: 4.0~6.0wt%, Al: 2.0~3.0wt%, W: 0.5~3.0wt%, Mo: 2.0~4.0wt%, C: 0.01~0.08wt%, Nb≤2.0wt%, Ta≤2.0wt%, Zr≤0.06wt%, B≤0.05wt%, and Ni is the balance. Preferably, in the nickel-based wrought superalloy: the sum of the weight percentages of Al, Ti, Nb, and Ta is greater than or equal to 7.0wt% and less than or equal to 9.5wt%; the sum of Nb and Ta is less than or equal to 2wt%; and the sum of W and Mo is greater than or equal to 3% and less than or equal to 6wt%. The design of the aforementioned chemical composition is explained as follows: This invention adds Al, Ti, Nb, and Ta elements to a nickel-based wrought superalloy, controlling the content of the γ' phase to 30-45 wt% (7.0% ≤ Al + Ti + Nb + Ta ≤ 9.5%, Nb + Ta ≤ 2%), thereby increasing the matrix channel width between the nanoscale γ' phases within the grains. Al, Ti, Nb, and Ta are all γ' phase forming elements, and the γ' phase is the most critical microstructure feature for ensuring the strength and creep life of the alloy. Therefore, the γ' phase content in the alloy is designed to reach 30-45 wt%. If the γ' phase content is further increased, it not only increases the difficulty of hot working of the alloy, but also narrows the matrix channel between two adjacent γ' phases, which will deteriorate the long-term microstructure stability of the alloy. Further control of 3% ≤ W + Mo ≤ 6% ensures that the alloy produces sufficient solid solution strengthening effect. Furthermore, adding boron (B) and boron (C) to the alloy to form carbides and borides can inhibit abnormal grain growth and strengthen the alloy. However, reducing the amount of carbides and borides prevents the formation of continuous thin-film grain boundary precipitates during long-term heat exposure. It should be noted that grain boundary slip is a key mechanism during high-temperature creep deformation. Therefore, strengthening grain boundaries with carborides can improve the alloy's creep life. However, excessive C and B content can cause continuous grain boundary precipitates during long-term service. Therefore, C is controlled at 0.01~0.08 wt%, and B ≤0.05 wt%. Cr is the main carbide-forming element and also improves the alloy's oxidation and corrosion resistance. However, excessive Cr content can lead to the formation of chromium precipitates (TCP), which is detrimental to the long-term stability of the alloy. Therefore, this invention controls the Cr content to 12~15%. In summary, through the above chemical composition design, nickel-based wrought superalloys exhibit excellent strength and long service life at 800~850℃, while also possessing good high-temperature machinability, thus enabling the production of nickel-based wrought superalloy parts with excellent strength and long service life.

[0028] On the other hand, embodiments of the present invention provide a method for preparing nickel-based deformable high-temperature alloy parts. First, vacuum induction melting, electroslag remelting, and vacuum arc remelting are used to prepare pure ingots conforming to the aforementioned chemical composition (in the above processes, the vacuum induction melted ingots and electroslag remelted ingots need to undergo stress-relief annealing to prevent excessive stress from causing cracking). The ingots are then subjected to diffusion annealing to eliminate segregation and improve the hot working properties of the alloy. The diffusion-annealed ingots are then subjected to multiple forging cycles at T-(10~40)℃~T+(10~30)℃ (T is the complete dissolution temperature of the γ´ phase of the alloy). The first 5~10 cycles are forged in the single-phase region to reduce deformation resistance and promote recrystallization, while the subsequent 5~10 cycles are forged in the two-phase region, producing a bar with a uniform microstructure. The bar is then loaded and unloaded, heated to T-(10~40)℃, and shaped to produce the formed part. The formed part is then subjected to solution treatment at T-20℃ to T+10℃ for 2 to 6 hours, followed by air cooling to room temperature. The purpose of solution treatment is to fully dissolve the γ´ phase and control the grain size. Subsequently, a first aging treatment is performed at 1000 to 1100℃ for 2 to 6 hours, followed by air cooling to room temperature, to induce a bimodal distribution of the γ´ phase size in the alloy. Finally, a second aging heat treatment is performed at 800 to 860℃ for 10 to 24 hours, followed by air cooling to room temperature to stabilize the γ´ phase, resulting in a nickel-based wrought superalloy part. In this invention, the chemical composition of the nickel-based wrought superalloy is designed to match the corresponding solution and aging treatment temperatures, resulting in nickel-based wrought superalloy parts with excellent strength and long service life.

[0029] 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

[0030] Figure 1 The grain structure of the nickel-based wrought superalloy hot-end component prepared in this embodiment;

[0031] Figure 2 The intragranular phase morphology of the nickel-based wrought superalloy hot-end component prepared in this embodiment;

[0032] Figure 3 The grain boundary precipitate morphology of the nickel-based wrought superalloy hot-end component prepared in this embodiment;

[0033] Figure 4 The grain structure of the hot-end component of the nickel-based wrought superalloy prepared in Comparative Example 1;

[0034] Figure 5 The intragranular phase morphology of the hot-end component of the nickel-based wrought superalloy prepared in Comparative Example 1;

[0035] Figure 6 The grain boundary precipitate morphology of the hot-end component of the nickel-based wrought superalloy prepared in Comparative Example 1 is shown. Detailed Implementation

[0036] 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 "embodiments" or "embodiments" 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.

[0037] To design a high-strength, long-life alloy that can withstand 800℃, this invention strengthens the alloy's properties by introducing low-energy layering fault energy design and introducing low-energy interface (microtwins) to replace part of the γ' phase (it should be noted that the low layering fault energy is related to the Co element content, and then the partial replacement of the γ' phase, the γ' phase content is related to the Al+Ti+Nb+Ta content, and the alloy contains 30-45wt% γ' phase, so a large number of microtwins will appear in the microstructure during tensile and long-term deformation); at the same time, the "low-carbon" design stabilizes the carbides and achieves long-term grain boundary stability.

[0038] In view of the above, the present invention achieves the design of a long-life alloy by reducing the effects of precipitation, solid solution and grain boundary strengthening through microtwin strengthening, so that the alloy can not only obtain high strength at 800℃, but also obtain excellent long service life.

[0039] This invention provides a nickel-based wrought superalloy, nickel-based wrought superalloy parts, and a method for preparing the same. The alloy not only exhibits excellent creep rupture at 800-850℃, but also good high-temperature machinability, enabling the fabrication of core hot-end components in novel engines and generators. This is achieved through the following technical solution:

[0040] On one hand, embodiments of the present invention provide a nickel-based wrought superalloy, wherein the chemical composition of the nickel-based wrought superalloy, by weight percentage, is as follows:

[0041] Co: 20~30wt%, Cr: 12~15wt%, Ti: 4.0~6.0wt%, Al: 2.0~3.0wt%, W: 0.5~3.0wt%, Mo: 2.0~4.0wt%, C: 0.01~0.08wt%, Nb≤2.0wt%, Ta≤2.0wt%, Zr≤0.06wt%, B≤0.05wt%, Ni is the balance.

[0042] Preferably, the chemical composition of the nickel-based wrought superalloy, by weight percentage, is as follows:

[0043] Co: 22.5~26.5wt%, Cr: 13.5~15.5wt%, Ti: 4.25~6.0wt%, Al: 2.2~2.6wt%, W: 1.1~1.5wt%, Mo: 2.5~3.5wt%, C: 0.01~0.04wt%, Nb: 0.5~1.0wt%, Ta: 0.5~1.0wt%, Zr: 0.025~0.04wt%, B: 0.01~0.04wt%, Ni is the balance.

[0044] Preferably, the chemical composition of the nickel-based wrought superalloy, by weight percentage, is as follows:

[0045] Co: 25wt%, Cr: 14wt%, Ti: 5.8wt%, Al: 2.4wt%, W: 1.2wt%, Mo: 2.8wt%, C: 0.04wt%, Nb: 1.0wt%, Ta: 1.0wt%, Zr: 0.015wt%, B: 0.015wt%, Ni is the balance.

[0046] Preferably, in the nickel-based wrought superalloy: the sum of the weight percentages of Al, Ti, Nb, and Ta is greater than or equal to 7.0 wt% and less than or equal to 9.5 wt%. The sum of Nb and Ta is less than or equal to 2 wt%; and the sum of W and Mo is greater than or equal to 3% and less than or equal to 6 wt%.

[0047] On the other hand, embodiments of the present invention provide a nickel-based wrought high-temperature alloy part, wherein the material of the nickel-based wrought high-temperature alloy part is any of the nickel-based wrought high-temperature alloys described above; preferably, the nickel-based wrought high-temperature alloy part is a hot-end component.

[0048] Preferably, the γ' phase content in the microstructure of the nickel-based deformable superalloy part is 30-45 wt%, and the γ' phase exhibits a bimodal distribution with peaks of 200-500 nm and <100 nm. More preferably, the morphology of the γ' phase is spherical particles. It should be noted that the bimodal intragranular γ' phase consists of discretely distributed large-sized secondary γ' phases and small-sized tertiary γ' phases distributed in the interstices of the secondary γ' phases. During creep, the secondary γ' phase coarsens by absorbing the surrounding tertiary γ' phase, and the tertiary γ' phases also aggregate and grow, gradually coarsening. The coarsening process of the secondary γ' phase can be divided into three stages: 1) initially discretely distributed spherical γ' phases; 2) gradual growth and coarsening by absorbing the surrounding tertiary γ' phase, with the morphology transforming towards a near-square shape; 3) tightly arranged square γ' phases, at which point the tertiary γ' phases are almost completely consumed, and the coarsening of the secondary γ' phase can only be completed through diffusion from surrounding Al and Ti elements, thus its size is relatively stable. Conversely, if the intracrystalline γ′ phase exhibits a unimodal distribution, it is composed of closely packed spherical γ′ phases. After a brief period of coarsening, the already narrow γ′ phase spacing becomes even smaller, thereby promoting the rafting of the initial alloy γ′ phase.

[0049] The average grain size of the nickel-based wrought superalloy part is grade 4-6; and / or the yield strength of the nickel-based wrought superalloy part at 800℃ is ≥800MPa, the tensile strength is ≥900MPa, and the elongation is ≥5%.

[0050] The nickel-based deformable high-temperature alloy parts have a creep rupture life of ≥2000h under tensile stress conditions of 800℃ and 150MPa.

[0051] The nickel-based deformable high-temperature alloy parts have a creep rupture life of ≥2000h under tensile stress conditions of 850℃ and 100MPa.

[0052] Furthermore, embodiments of the present invention provide a method for preparing the above-mentioned nickel-based deformed high-temperature alloy parts, which includes the following steps:

[0053] Step 1): The raw materials are subjected to vacuum induction melting, electroslag remelting, and vacuum consumable melting to obtain nickel-based deformed high-temperature alloy ingots.

[0054] In this step, the raw material is subjected to vacuum induction melting to obtain a vacuum induction melt ingot; the vacuum induction melt ingot is first subjected to a first stress-relief annealing treatment followed by an electroslag remelting treatment; the electroslag remelted ingot obtained after the electroslag remelting treatment is first subjected to a second stress-relief annealing treatment followed by a vacuum arc remelting treatment to obtain a nickel-based wrought superalloy ingot. Preferably, the temperature of the first stress-relief annealing treatment is 900-1050℃, and the time of the first stress-relief annealing treatment is 5-15h; preferably, the temperature of the second stress-relief annealing treatment is 900-1050℃, and the time of the second stress-relief annealing treatment is 5-15h.

[0055] Step 2): The nickel-based deformed superalloy ingot is subjected to diffusion annealing to obtain the diffusion-annealed ingot.

[0056] The nickel-based deformed high-temperature alloy ingot is heated to 1140-1170℃ and held for 20-50 hours, then heated to 1170-1200℃ and held for 20-50 hours. After being cooled to 600℃ in the furnace, it is taken out of the furnace to obtain the ingot after diffusion annealing.

[0057] Step 3): The ingot after diffusion annealing is subjected to multi-fire forging to obtain nickel-based deformed high-temperature alloy bars;

[0058] In this step, the diffusion-annealed ingot is subjected to multiple forging processes at a temperature of T - (10~40) °C to T + (10~30) °C, where T is the complete dissolution temperature of the γ´ phase in the nickel-based wrought superalloy. Preferably, the first 5~10 forging processes are performed in the single-phase region to reduce deformation resistance and promote recrystallization, and the last 5~10 forging processes are performed in the two-phase region. The temperature of the single-phase region is T to T + (10~30) °C; the temperature of the two-phase region is T - (10~40) °C to T.

[0059] Step 4): The nickel-based deformed high-temperature alloy bar is subjected to forming treatment (e.g., die forging and extrusion forming) to obtain a formed part.

[0060] The forming process temperature is T - (10~40)℃; where T is the complete dissolution temperature of the γ´ phase of the nickel-based wrought superalloy.

[0061] Step 5): Perform solution treatment on the molded part to obtain the solution-treated molded part.

[0062] In step 5), the solution treatment temperature is (T-20)℃ to (T+10)℃, and the solution treatment time is 2-6 hours. Here, T is the complete dissolution temperature of the γ´ phase of the nickel-based wrought superalloy.

[0063] Step 6): The solution-treated formed part is subjected to aging treatment to obtain a nickel-based deformed high-temperature alloy part.

[0064] In step 6): the solution-treated molded part is heated to 1000-1100℃ and held for 2-6 hours for a first aging treatment. After cooling, the molded part after the first aging treatment is obtained. The molded part after the first aging treatment is heated to 800-860℃ and held for 10-24 hours for a second aging treatment. After cooling, the nickel-based deformed high-temperature alloy part is obtained.

[0065] In summary, the embodiments of this invention provide a nickel-based wrought superalloy, nickel-based wrought superalloy parts, and their preparation method. The designed nickel-based wrought superalloy is a high-temperature, long-life wrought superalloy. The content of precipitated phase (γ´ phase) in the alloy is about 10% higher than that of conventional ultra-supercritical wrought superalloys, but significantly lower than that of alloys used in aero-engines that withstand 800℃ and above. The alloy matrix channels are significantly widened, while the amount of carbides and borides at grain boundaries is reduced. The alloy is strengthened by low-energy interfaces such as microtwins and stacking faults induced by low stacking fault energy, and the grain size is controlled by heat treatment to reduce the grain boundary content, thereby achieving a coordinated strengthening effect within the grains and significantly improving the long-term structural stability of the alloy.

[0066] The following detailed description uses specific examples.

[0067] Table 1 shows the chemical composition (wt.%) of the alloys in the embodiments and comparative examples of the present invention.

[0068] Table 1

[0069]

[0070] Example 1

[0071] This embodiment describes the preparation of a nickel-based wrought high-temperature alloy part (hot-end component), including the following steps:

[0072] Step 1): A pure nickel-based wrought superalloy ingot (508 mm in diameter) is prepared using vacuum induction melting, electroslag remelting, and vacuum arc remelting. During these processes, the vacuum induction melted ingot and the electroslag remelted ingot undergo stress-relief annealing (at a temperature of 930°C for 10 hours) to prevent cracking caused by excessive stress.

[0073] Step 2): The nickel-based wrought superalloy ingot is subjected to diffusion annealing to obtain a diffusion-annealed ingot. Specifically, the nickel-based wrought superalloy ingot is heated to 1140°C, held for 25 hours, then heated to 1180°C, held for 40 hours, and then cooled in the furnace to below 600°C to obtain the diffusion-annealed ingot.

[0074] Step 3): The diffusion-annealed ingot is forged 18 times in the temperature range of 1090-1150 ℃. The first 5 forging times are at 1150 ℃, the middle 6 forging times are at 1130 ℃, and the last 7 forging times are at 1090 ℃, to prepare a nickel-based wrought superalloy bar with a diameter of 250 mm and uniform microstructure. The complete dissolution temperature T of the γ´ phase of the nickel-based wrought superalloy is 1130 ℃.

[0075] Step 4): Then, the nickel-based deformed high-temperature alloy bar is loaded and unloaded, heated to 1120℃ to form the hot-end component, and the formed part is obtained.

[0076] Step 5): The molded part is subjected to solution treatment at 1130℃ for 4 hours to allow the γ´ phase to fully dissolve back, and then air-cooled to room temperature to obtain the solution-treated molded part.

[0077] Step 6): Heat the solution-treated molded part to 1050℃ and hold for 4 hours for the first aging treatment, air cool to room temperature, then heat to 850℃ and hold for 24 hours for the second aging treatment, and air cool to room temperature to obtain a nickel-based deformed high-temperature alloy hot end part.

[0078] in, Figure 1 The grain structure of the nickel-based wrought superalloy hot-end component prepared in this embodiment; Figure 2 The intragranular phase morphology of the nickel-based wrought superalloy hot-end component prepared in this embodiment; Figure 3 This image shows the grain boundary precipitate morphology of the nickel-based wrought superalloy hot-end component prepared in this embodiment. Specifically, from... Figure 1 It can be seen that in the nickel-based deformable superalloy hot-end component prepared in this embodiment, the γ′ phase content in the alloy microstructure is 30-45 wt%, and there is basically no primary γ′ phase at the grain boundaries. Therefore, the γ′ phase is all inside the grains, and the average grain size of the alloy is grade 5. From Figure 2 It can be seen that in the nickel-based wrought superalloy hot-end component prepared in this embodiment, the γ′ phase in the grains exhibits a bimodal distribution, namely a large-size γ′ phase of 200~500nm and a small-size γ′ phase of <100nm. From Figure 3 It can be seen that there are a small amount of granular carbides on the grain boundaries, and the grain boundaries have a serrated morphology.

[0079] The performance of the nickel-based wrought superalloy hot-end component prepared in this embodiment was tested. The nickel-based wrought superalloy prepared in this embodiment exhibited a yield strength of 880 MPa, a tensile strength of 1090 MPa, and an elongation of 6.5% at 800°C. The crease service life of the nickel-based wrought superalloy prepared in this embodiment was 3600 h at 800°C / 150 MPa. The crease service life of the nickel-based wrought superalloy prepared in this embodiment was 3200 h at 850°C / 100 MPa. The performance testing standards for this embodiment and the following embodiments are as follows: the tensile property testing standard at 800°C is GB / T 228.2; the crease service life testing standard is GB / T 2039.

[0080] Example 2

[0081] This embodiment describes the preparation of a nickel-based wrought high-temperature alloy part (hot-end component), including the following steps:

[0082] Step 1): A pure nickel-based wrought superalloy ingot (508 mm in diameter) is prepared using vacuum induction melting, electroslag remelting, and vacuum arc remelting. During these processes, the vacuum induction melted ingot and the electroslag remelted ingot undergo stress-relief annealing (at a temperature of 980°C for 8 hours) to prevent cracking caused by excessive stress.

[0083] Step 2): The nickel-based wrought superalloy ingot is subjected to diffusion annealing to obtain a diffusion-annealed ingot. Specifically, the nickel-based wrought superalloy ingot is heated to 1150°C, held for 35 hours, then heated to 1190°C, held for 40 hours, and then cooled in the furnace to below 600°C to obtain the diffusion-annealed ingot.

[0084] Step 3): The diffusion-annealed ingot is forged 25 times in the temperature range of 1130-1170 ℃. The first 5 forging times are at 1170 ℃, the middle 10 times are at 1160 ℃, and the last 10 times are at 1130 ℃. This produces a nickel-based wrought superalloy bar with a diameter of 300 mm and uniform microstructure. The complete phase reversion temperature T of the nickel-based wrought superalloy is 1140 ℃.

[0085] Step 4): Then, the nickel-based deformed high-temperature alloy bar is loaded and unloaded, heated to 1130℃ to form the hot end component, and the formed part is obtained.

[0086] Step 5): The molded part is subjected to solution treatment at 1140℃ for 4 hours to allow the γ´ phase to fully dissolve back, and then air-cooled to room temperature to obtain the solution-treated molded part.

[0087] Step 6): Heat the solution-treated molded part to 1050℃ and hold for 4 hours for the first aging treatment, air cool to room temperature, then heat to 850℃ and hold for 24 hours for the second aging treatment, and air cool to room temperature to obtain a nickel-based deformed high-temperature alloy hot end part.

[0088] The performance of the nickel-based wrought superalloy hot-end component prepared in this embodiment was tested. The nickel-based wrought superalloy prepared in this embodiment exhibited a yield strength of 870 MPa, a tensile strength of 1050 MPa, and an elongation of 8% at 800°C. The crease service life of the nickel-based wrought superalloy prepared in this embodiment was 3900 h at 800°C / 150 MPa. The crease service life of the nickel-based wrought superalloy prepared in this embodiment was 3500 h at 850°C / 100 MPa.

[0089] In the nickel-based wrought superalloy hot-end component prepared in this embodiment, the γ′ phase content in the alloy microstructure is 30-45 wt%, with virtually no primary γ′ phase at grain boundaries. Therefore, the γ′ phase is entirely within the grains, and the average grain size of the alloy is grade 5. In the nickel-based wrought superalloy hot-end component prepared in this embodiment, the γ′ phase in the grains exhibits a bimodal distribution, consisting of a large-size γ′ phase of 200-500 nm and a small-size γ′ phase of <100 nm. A small amount of granular carbides are present at the grain boundaries, and the grain boundaries exhibit a serrated morphology.

[0090] Example 3

[0091] This embodiment describes the preparation of a nickel-based wrought high-temperature alloy part (hot-end component), including the following steps:

[0092] Step 1): A pure nickel-based wrought superalloy ingot (508 mm in diameter) is prepared using vacuum induction melting, electroslag remelting, and vacuum arc remelting. During these processes, the vacuum induction melted ingot and the electroslag remelted ingot undergo stress-relief annealing (at a temperature of 1000°C for 10 hours) to prevent cracking caused by excessive stress.

[0093] Step 2): The nickel-based wrought superalloy ingot is subjected to diffusion annealing to obtain a diffusion-annealed ingot. Specifically, the nickel-based wrought superalloy ingot is heated to 1160°C, held for 40 hours, then heated to 1190°C, held for 50 hours, and then cooled in the furnace to below 600°C to obtain the diffusion-annealed ingot.

[0094] Step 3): The diffusion-annealed ingot is forged 20 times in the temperature range of 1140-1180 ℃. The first 5 forging times are at 1170 ℃, the middle 5 times are at 1160 ℃, and the last 5 times are at 1140 ℃. This produces a nickel-based wrought superalloy bar with a diameter of 200 mm and uniform microstructure. The complete phase reversion temperature T of the nickel-based wrought superalloy is 1160 ℃.

[0095] Step 4): Then, the nickel-based deformed high-temperature alloy bar is loaded and unloaded, heated to 1130℃ to form the hot end component, and the formed part is obtained.

[0096] Step 5): The molded part is subjected to solution treatment at 1145℃ for 4 hours to allow the γ´ phase to fully dissolve back, and then air-cooled to room temperature to obtain the solution-treated molded part.

[0097] Step 6): Heat the solution-treated molded part to 1030℃ and hold for 6 hours for the first aging treatment, air cool to room temperature, then heat to 850℃ and hold for 24 hours for the second aging treatment, and air cool to room temperature to obtain a nickel-based deformed high-temperature alloy hot end part.

[0098] The performance of the nickel-based wrought superalloy hot-end component prepared in this embodiment was tested. The nickel-based wrought superalloy prepared in this embodiment exhibited a yield strength of 850 MPa, a tensile strength of 1060 MPa, and an elongation of 9% at 800°C. The crease service life of the nickel-based wrought superalloy prepared in this embodiment was 3600 h at 800°C / 150 MPa. The crease service life of the nickel-based wrought superalloy prepared in this embodiment was 3100 h at 850°C / 100 MPa.

[0099] In the nickel-based wrought superalloy hot-end component prepared in this embodiment, the γ′ phase content in the alloy microstructure is 30-45 wt%, with virtually no primary γ′ phase at grain boundaries. Therefore, the γ′ phase is entirely within the grains, and the average grain size of the alloy is grade 5. In the nickel-based wrought superalloy hot-end component prepared in this embodiment, the γ′ phase in the grains exhibits a bimodal distribution, consisting of a large-size γ′ phase of 200-500 nm and a small-size γ′ phase of <100 nm. A small amount of granular carbides are present at the grain boundaries, and the grain boundaries exhibit a serrated morphology.

[0100] Comparative Example 1

[0101] This embodiment describes the preparation of a nickel-based wrought high-temperature alloy part (hot-end component), including the following steps:

[0102] Step 1): A pure nickel-based wrought superalloy ingot (508 mm in diameter) was prepared by vacuum induction melting, electroslag remelting and vacuum arc remelting.

[0103] Step 2): The nickel-based wrought superalloy ingot is subjected to diffusion annealing to obtain a diffusion-annealed ingot. Specifically, the nickel-based wrought superalloy ingot is heated to 1160°C, held at that temperature for 50 hours, then heated to 1190°C, held at that temperature for 50 hours, and then cooled in the furnace to below 600°C to obtain the diffusion-annealed ingot.

[0104] Step 3): The diffusion-annealed ingot is forged 20 times in the temperature range of 1130-1170 ℃ to prepare a bar with a diameter of 200 mm and uniform microstructure.

[0105] Step 4): Load and unload the bar stock, heat it to 1140 ℃ to form the hot end component, and obtain the formed part.

[0106] Step 5): The molded part is solution treated at 1145℃ for 2 hours and then air-cooled to room temperature; then it is kept at 850℃ for 4 hours and air-cooled to room temperature, and then kept at 780℃ for 16 hours and air-cooled to room temperature to obtain the hot end part.

[0107] Figure 4The grain structure of the hot-end component of the nickel-based wrought superalloy prepared in Comparative Example 1; Figure 5 The intragranular phase morphology of the hot-end component of the nickel-based wrought superalloy prepared in Comparative Example 1; Figure 6 The image shows the grain boundary precipitate morphology of the nickel-based wrought superalloy hot-end component prepared in Comparative Example 1. Among them, from... Figure 4 It can be seen that the alloy contains micron-sized primary γ′ phases at the grain boundaries, and the average grain size of the alloy is grade 8. From... Figure 5 It can be seen that the γ′ phase in the grains exhibits a unimodal distribution, with only large-sized γ′ phases on the order of hundreds of nanometers. From... Figure 6 It can be seen that there are a lot of carbides on the grain boundaries, and some of them are already connected to each other, and the grain boundaries are relatively straight.

[0108] Performance was tested on the hot-end component of Comparative Example 1. The hot-end component of Comparative Example 1 had a yield strength of 860 MPa, a tensile strength of 1080 MPa, and an elongation of 5% at 800℃. The creep rupture life was 500 h at 800℃ / 150 MPa and 400 h at 850℃ / 100 MPa.

[0109] The reason for the poor performance of the hot-end component prepared in Comparative Example 1 is that the micron-sized γ′ phase at the grain boundary does not play a strengthening role, the γ′ phase in the grain is distributed in a single peak, the distance between adjacent γ′ phases is short, and it is easy to coarsen or raft. There are many grain boundary carbides, which are easy to form continuous carbides during long-term service. Therefore, although the alloy has high strength, its service life is low.

[0110] It should be noted that in Comparative Example 1, because the combined content of Al, Ti, and Nb exceeds 9.5 wt%, the alloy has a high content of the γ′ phase, resulting in short distances between adjacent γ′ phases and a tendency for coarsening or rafting. Furthermore, the C content reaches 0.08 wt%, leading to abundant grain boundary carbides, which easily form continuous carbides during long-term service. These two factors contribute to the alloy's relatively low long-term creep life.

[0111] Comparative Example 2

[0112] This paper describes the comparative preparation of a nickel-based wrought superalloy part (hot-end component), including the following steps:

[0113] Step 1): A pure nickel-based wrought superalloy ingot (508 mm in diameter) is prepared using vacuum induction melting, electroslag remelting, and vacuum arc remelting. During these processes, the vacuum induction melted ingot and the electroslag remelted ingot undergo stress-relief annealing (at a temperature of 1050°C for 10 hours) to prevent cracking caused by excessive stress.

[0114] Step 2): The nickel-based wrought superalloy ingot is subjected to diffusion annealing to obtain a diffusion-annealed ingot. Specifically, the nickel-based wrought superalloy ingot is heated to 1160°C, held for 50 hours, then heated to 1190°C, held for 50 hours, and then cooled in the furnace to below 600°C to obtain a diffusion-annealed ingot.

[0115] Step 3): The diffusion-annealed ingot is forged 25 times in the temperature range of 1140-1190 ℃. The first 5 forging times are at 1190 ℃, the middle 15 forging times are at 1160 ℃, and the last 5 forging times are at 1140 ℃. This produces a nickel-based wrought superalloy bar with a diameter of 200 mm and uniform microstructure. The complete phase reversion temperature T of the nickel-based wrought superalloy is 1180 ℃.

[0116] Step 4): Then, the nickel-based deformed high-temperature alloy bar is loaded and unloaded, heated to 1140℃ to form the hot-end component, and the formed part is obtained.

[0117] Step 5): The molded part is subjected to solution treatment at 1150℃ for 4 hours to allow the γ´ phase to fully dissolve back, and then air-cooled to room temperature to obtain the solution-treated molded part.

[0118] Step 6): Heat the solution-treated formed part to 850℃ and hold for 24 hours for aging treatment, then air cool to room temperature to obtain a nickel-based deformed high-temperature alloy hot end part.

[0119] Performance tests were conducted on hot-end components of nickel-based wrought superalloys prepared in the comparative example. The nickel-based wrought superalloy prepared in this comparative example exhibited a tensile strength of 1150 MPa, a yield strength of 800 MPa, and an elongation of 8% at 800°C. The nickel-based wrought superalloy prepared in this example had a creep rupture life of 760 h at 800°C / 150 MPa. The nickel-based wrought superalloy prepared in this example had a creep rupture life of 540 h at 850°C / 100 MPa.

[0120] Because the alloy in Comparative Example 2 uses traditional single-step aging, the γ′ phase in the alloy has a single-peak distribution. The adjacent γ′ phase matrix channels are narrow and easily coarsened or rafted, resulting in a low creep life.

[0121] It should be noted that the total content of the γ′ phase forming elements Al, Ti, Nb and Ta in Comparative Example 2 is relatively high, so the γ′ phase content in the alloy can reach 60%. Although this can ensure high-temperature strength, long-term structural stability cannot be guaranteed, so it is prone to coarsening or rafting.

[0122] Comparative Example 3

[0123] The alloy composition of this comparative example is the same as that of Example 2. A nickel-based wrought high-temperature alloy part (hot-end component) is prepared by the following steps:

[0124] Step 1): A pure nickel-based wrought superalloy ingot (508 mm in diameter) is prepared using vacuum induction melting, electroslag remelting, and vacuum arc remelting. During these processes, the vacuum induction melted ingot and the electroslag remelted ingot undergo stress-relief annealing (at a temperature of 980°C for 8 hours) to prevent cracking caused by excessive stress.

[0125] Step 2): The nickel-based wrought superalloy ingot is subjected to diffusion annealing to obtain a diffusion-annealed ingot. Specifically, the nickel-based wrought superalloy ingot is heated to 1160°C, held at that temperature for 50 hours, then heated to 1190°C, held at that temperature for 50 hours, and cooled to obtain a diffusion-annealed ingot.

[0126] Step 3): The diffusion-annealed ingot is forged 15 times at 1180℃ and 10 times at 1120℃. This produces a nickel-based wrought superalloy bar with a diameter of 200mm and uniform microstructure.

[0127] Step 4): Then, the nickel-based deformed high-temperature alloy bar is loaded and unloaded, heated to 1120℃ to form the hot-end component, and the formed part is obtained.

[0128] Step 5): The molded part is subjected to solution treatment at 1120℃ for 4 hours, and then air-cooled to room temperature to obtain the solution-treated molded part.

[0129] Step 6): Heat the solution-treated formed part to 850℃ and hold for 24 hours for aging treatment, then air cool to room temperature to obtain a nickel-based deformed high-temperature alloy hot end part.

[0130] Performance was tested on the hot-end component of the nickel-based wrought superalloy prepared in Comparative Example 3. The nickel-based wrought superalloy prepared in this comparative example exhibited a yield strength of 750 MPa, a tensile strength of 900 MPa, and an elongation of 8% at 800°C. The nickel-based wrought superalloy prepared in this example had a creep rupture life of 600 h at 800°C / 150 MPa. The nickel-based wrought superalloy prepared in this example had a creep rupture life of 350 h at 850°C / 100 MPa.

[0131] Here, because the deformation temperature of Comparative Example 3 is relatively low (i.e., the forging temperature of the final firing and the preparation temperature of the formed part are relatively low), the final grain size is relatively fine. Then, no two-step aging treatment is performed, so the γ′ phase has a single-peak distribution, resulting in relatively low high-temperature strength and creep life.

[0132] 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 nickel-based wrought superalloy, characterized in that, The chemical composition of the nickel-based wrought superalloy, by weight percentage, is as follows: Co: 20~30wt%, Cr: 12~15wt%, Ti: 4.0~6.0wt%, Al: 2.0~3.0wt%, W: 0.5~3.0wt%, Mo: 2.0~4.0wt%, C: 0.01~0.08wt%, Nb≤2.0wt%, Ta≤2.0wt%, Zr≤0.06wt%, B≤0.05wt%, Ni is the balance.

2. The nickel-based wrought superalloy according to claim 1, characterized in that, The chemical composition of the nickel-based wrought superalloy, by weight percentage, is as follows: Co: 22.5~26.5wt%, Cr: 13.5~15.5wt%, Ti: 4.25~6.0wt%, Al: 2.2~2.6wt%, W: 1.1~1.5wt%, Mo: 2.5~3.5wt%, C: 0.01~0.04wt%, Nb: 0.5~1.0wt%, Ta: 0.5~1.0wt%, Zr: 0.025~0.04wt%, B: 0.01~0.04wt%, Ni is the balance.

3. The nickel-based wrought superalloy according to claim 1 or 2, characterized in that, In the nickel-based wrought superalloy: The sum of the weight percentages of Al, Ti, Nb, and Ta elements is greater than or equal to 7.0 wt% and less than or equal to 9.5 wt%; and / or The sum of Nb and Ta elements is less than or equal to 2wt%; and / or The sum of W and Mo elements is greater than or equal to 3% and less than or equal to 6 wt%.

4. A nickel-based wrought high-temperature alloy part, characterized in that, The material of the nickel-based wrought superalloy part is the nickel-based wrought superalloy as described in any one of claims 1-3; Preferably, the nickel-based deformed high-temperature alloy part is a hot-end component.

5. The nickel-based wrought high-temperature alloy part according to claim 4, characterized in that, The nickel-based deformable superalloy part has a γ' phase content of 30-45 wt% in its microstructure, and the γ' phase exhibits a bimodal distribution at 200-500 nm and <100 nm; and / or The morphology of the γ´ phase in the alloy microstructure of the nickel-based deformed superalloy part is spherical particles; and / or The average grain size of the nickel-based deformable superalloy part is grade 4-6; and / or The nickel-based deformable superalloy parts exhibit a yield strength ≥800MPa, tensile strength ≥900MPa, and elongation ≥5% at 800℃; and / or The nickel-based deformable superalloy part exhibits a creep rupture life ≥2000h under tensile stress conditions of 800℃ and 150MPa; and / or The nickel-based deformable high-temperature alloy parts have a creep rupture life of ≥2000h under tensile stress conditions of 850℃ and 100MPa.

6. The method for preparing nickel-based wrought superalloy parts according to claim 4 or 5, characterized in that, It includes the following steps: Step 1): The raw materials are subjected to vacuum induction melting, electroslag remelting, and vacuum consumable melting to obtain nickel-based deformed high-temperature alloy ingots; Step 2): The nickel-based deformed superalloy ingot is subjected to diffusion annealing to obtain a diffusion-annealed ingot; Step 3): The ingot after diffusion annealing is subjected to multi-fire forging to obtain nickel-based deformed high-temperature alloy bars; Step 4): The nickel-based wrought high-temperature alloy bar is formed to obtain a shaped part; Step 5): Perform solution treatment on the molded part to obtain the solution-treated molded part; Step 6): The solution-treated formed part is subjected to aging treatment to obtain a nickel-based deformed high-temperature alloy part.

7. The method for preparing nickel-based wrought superalloy parts according to claim 6, characterized in that, In step 1): The raw materials are subjected to vacuum induction melting to obtain a vacuum induction melting ingot; the vacuum induction melting ingot is first subjected to a first stress-relief annealing treatment and then subjected to an electroslag remelting treatment; the electroslag remelting ingot obtained after the electroslag remelting treatment is first subjected to a second stress-relief annealing treatment and then subjected to vacuum arc remelting treatment to obtain a nickel-based deformed high-temperature alloy ingot. Preferably, the temperature of the first stress-relieving annealing treatment is 900-1050℃, and the time of the first stress-relieving annealing treatment is 5-15h; Preferably, the temperature of the second stress-relieving annealing treatment is 900-1050℃, and the time of the second stress-relieving annealing treatment is 5-15h.

8. The method for preparing nickel-based wrought superalloy parts according to claim 6, characterized in that, In step 2): The nickel-based wrought superalloy ingot is heated to 1140-1170℃ and held for 20-50 hours. Then it is heated to 1170-1200℃ and held for 20-50 hours. After being cooled in the furnace to below 600℃, the ingot is obtained after diffusion annealing.

9. The method for preparing nickel-based wrought superalloy parts according to claim 6, characterized in that, In step 3): The diffusion-annealed ingot is subjected to multiple forging processes at a temperature of T-(10~40) ℃ to T+(10~30) ℃; where T is the complete dissolution temperature of the γ´ phase of the nickel-based wrought superalloy. Preferably, the first 5 to 10 forging cycles are performed in the single-phase region to reduce deformation resistance and promote recrystallization, and the last 5 to 10 forging cycles are performed in the two-phase region; preferably, the temperature of the single-phase region is T to T+(10~30) ℃; the temperature of the two-phase region is T-(10~40) ℃ to T.

10. The method for preparing nickel-based wrought superalloy parts according to claim 6, characterized in that, In step 4): the forming treatment temperature is T - (10~40)℃; where T is the complete dissolution temperature of the γ´ phase of the nickel-based wrought superalloy; and / or In step 5): the solution treatment temperature is (T-20)℃~(T+10)℃, and the solution treatment time is 2-6h; where T is the complete dissolution temperature of the γ´ phase of the nickel-based wrought superalloy; and / or In step 6): the solution-treated molded part is heated to 1000-1100℃ and held for 2-6 hours for a first aging treatment. After cooling, the molded part after the first aging treatment is obtained. The molded part after the first aging treatment is heated to 800-860℃ and held for 10-24 hours for a second aging treatment. After cooling, the nickel-based deformed high-temperature alloy part is obtained.

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