A nickel-based superalloy and a method of making the same

Abstract

The present application relates to the alloy technical field, especially to a kind of nickel-based superalloy and preparation method thereof.The present application provides a kind of nickel-based superalloy, according to mass percentage, including the following components: C 0.04%~0.06%, Co 15.0%~15.04%, Cr 10.08%~10.13%, W 5.74%~5.81%, Mo 3.30%~3.34%, Al 3.73%~3.78%, Ti 3.68%~3.73%, Nb 1.77%~1.81%, Hf 0.20%~0.40%, B 0.010%~0.014%, Zr 0.008%~0.01%, Mg 0.002%~0.006%, the balance is Ni.The nickel-based superalloy of the present application, by adjusting component, compared with FGH4097 alloy has more excellent comprehensive performance.

Description

Technical Field

[0001] This invention relates to the field of alloy technology, and in particular to a nickel-based superalloy and its preparation method. Background Technology

[0002] FGH4097 alloy, as a precipitation-strengthened nickel-based powder superalloy, has a γ′ phase mass fraction of up to about 61%, exhibiting excellent performance in high-temperature environments. In particular, its comprehensive mechanical properties are outstanding in the temperature range of 650℃ to 750℃, making it an ideal material for manufacturing key hot-end components in high-end equipment such as aero-engines.

[0003] Direct hot isostatic pressing (HIP) is a near-net-shape forming technology for the fabrication of powder metallurgy superalloy discs. Components such as FGH4097 alloy discs, shafts, and rings prepared using this technology not only exhibit uniform microstructure but also demonstrate consistent performance, significantly improving product reliability and service life. This process reduces material waste, increases production efficiency, and ensures product quality, making it an ideal choice for manufacturing high-end powder metallurgy superalloy components.

[0004] Currently, FGH4097 alloy is widely used in key components such as turbine disks and compressor disks of aero-engines and ground-based gas turbines. These components are subjected to extremely high temperatures and pressures within the engine, thus requiring materials with excellent high-temperature, high-pressure, and fatigue resistance. With the continuous development of high thrust-to-weight ratio aero-engine technology, the operating temperature and performance requirements for key hot-end engine components such as turbine disks are constantly increasing. To meet the demands of higher-performance engines, further in-depth research and improvement of FGH4097 alloy are needed to further enhance its high-temperature, high-pressure, and fatigue resistance, providing stronger support for the development of high-end equipment such as aero-engines.

[0005] In view of this, this invention is hereby proposed. Summary of the Invention

[0006] The primary objective of this invention is to provide a nickel-based superalloy that, by adjusting the alloy composition, exhibits higher tensile strength and high-temperature durability compared to the FGH4097 alloy, demonstrating superior overall performance.

[0007] The second objective of this invention is to provide a method for preparing the above-mentioned nickel-based superalloy, which improves the tensile strength and high-temperature durability of the nickel-based superalloy.

[0008] In order to achieve the above-mentioned objectives of the present invention, the following technical solution is adopted:

[0009] In a first aspect, the present invention provides a nickel-based superalloy comprising, by weight percentage, the following components:

[0010] C 0.05%–0.063%, Co 15.0%–15.04%, Cr 10.08%–10.13%, W 5.74%–5.81%, Mo 3.30%–3.34%, Al 3.73%–3.78%, Ti 3.68%–3.73%, Nb 1.77%–1.81%, Hf 0.20%–0.40%, B 0.010%–0.014%, Zr 0.008%–0.01%, Mg 0.002%–0.006%, balance Ni.

[0011] Furthermore, in the nickel-based superalloy, the mass ratio of Ti to Al is (0.97–1):1.

[0012] Furthermore, in the nickel-based superalloy, the total mass percentage of Al, Ti, Nb, and Hf is 9.38% to 9.71%.

[0013] Furthermore, it includes at least one of the following features (1) to (4);

[0014] (1) The nickel-based superalloy includes γ phase, γ' phase, carbide, boride and TCP phase;

[0015] (2) The average grain size of the nickel-based superalloy is 25-35 μm;

[0016] (3) The mass percentage of the γ′ phase in the nickel-based superalloy is 57% to 59%;

[0017] (4) The average size of the primary γ′ phase of the nickel-based superalloy is 0.5-4 μm, the average size of the secondary γ′ phase is 0.2-0.8 μm, and the average size of the tertiary γ′ phase is 0.04-0.07 μm.

[0018] Furthermore, it includes at least one of the following features (1) to (3);

[0019] (1) The room temperature tensile strength of the nickel-based superalloy is 1510-1590 MPa, the room temperature yield strength is 1020-1185 MPa, the room temperature elongation after fracture is 14%-25%, and the room temperature reduction of area is 17%-23%.

[0020] (2) The nickel-based high-temperature alloy has a tensile strength of 1090-1250 MPa, a yield strength of 913-1050 MPa, an elongation after fracture of 12%-35%, and a reduction of area of ​​12%-32% at 750℃.

[0021] (3) The nickel-based high-temperature alloy has a composite creep rupture time of 180-220h at 650℃ and 990MPa and a composite creep rupture time of 190-210h at 750℃ and 637MPa.

[0022] Secondly, the present invention also provides a method for preparing the nickel-based superalloy as described above, comprising the following steps:

[0023] Alloy powder is prepared by mixing ingredients according to the specified ratio; the alloy powder is then subjected to hot isostatic pressing, solution treatment and aging treatment in sequence to obtain the nickel-based high-temperature alloy.

[0024] Furthermore, the hot isostatic pressing treatment includes: heat treatment at 1190-1220℃ and 130-150MPa for 3-5 hours.

[0025] Furthermore, the solution treatment includes: holding at 1170–1220°C for 2–5 hours, followed by cooling.

[0026] Furthermore, the cooling method for the solution treatment includes furnace cooling followed by air cooling; or, the cooling method for the solution treatment includes oil cooling.

[0027] And / or, the cooling rate of the solution treatment is 30-40°C / min; or, the cooling rate of the solution treatment is ≥80°C / min.

[0028] Furthermore, the aging treatment includes: heat treatment at 870±10℃ for 6 to 9 hours, followed by air cooling, heat treatment at 760±10℃ for 14 to 18 hours, followed by air cooling.

[0029] Compared with the prior art, the beneficial effects of the present invention are as follows:

[0030] 1. The nickel-based superalloy provided by this invention adjusts the alloy composition based on the FGH4097 alloy composition. By increasing the C, Cr, and W content and decreasing the Co and Mo content, the solid solution strengthening effect is improved. By increasing the Ti content and decreasing the Al and Nb content, the complete solid solution temperature of the γ′ phase is increased while keeping the total amount of γ′ phase forming elements Al, Ti, Nb, and Hf basically unchanged. As a result, its comprehensive properties such as tensile properties, creep rupture properties, impact properties, and fatigue properties are better than those of the FGH4097 alloy, especially its tensile strength and high-temperature creep rupture properties are superior to those of the FGH4097 alloy.

[0031] 2. The preparation method of the nickel-based superalloy of the present invention improves the performance of the nickel-based superalloy by studying the hot isostatic pressing and heat treatment processes. The room temperature tensile strength, high temperature tensile strength and creep strength of the prepared nickel-based superalloy are all better than those of the FGH4097 alloy prepared by direct hot isostatic pressing near-net-shape forming. Attached Figure Description

[0032] To more clearly illustrate the technical solutions of the embodiments of the present invention, the accompanying drawings used in the embodiments will be briefly introduced below. It should be understood that the following drawings only show some embodiments of the present invention and should not be regarded as a limitation of the scope. For those skilled in the art, other related drawings can be obtained from these drawings without creative effort.

[0033] Figure 1 This is the grain structure of the small ingot blank in Embodiment 1 of the present invention.

[0034] Figure 2 This is the grain structure of the small ingot blank in Embodiment 2 of the present invention.

[0035] Figure 3 This is the grain structure of the small ingot blank in Embodiment 3 of the present invention.

[0036] Figure 4 This is the grain structure of the small ingot blank in Example 4 of the present invention.

[0037] Figure 5 This is the γ′ phase characteristic of the small ingot blank in Embodiment 1 of the present invention.

[0038] Figure 6 This is the γ′ phase characteristic of the small ingot blank in Embodiment 2 of the present invention.

[0039] Figure 7 This is the γ′ phase characteristic of the small ingot blank in Embodiment 3 of the present invention.

[0040] Figure 8 This is the γ′ phase characteristic of the small ingot blank in Embodiment 4 of the present invention.

[0041] Figure 9 This is the grain structure of the DP-HT-OC-1 small ingot blank of the present invention.

[0042] Figure 10 This is the grain structure of the DP-HT-OC-3 small ingot blank of the present invention.

[0043] Figure 11 This is the grain structure of the DP-HT-OC-9 small ingot blank of the present invention.

[0044] Figure 12 This describes the γ′ phase characteristics of the DP-HT-OC-1 small ingot blank of the present invention.

[0045] Figure 13 This describes the γ′ phase characteristics of the DP-HT-OC-3 small ingot blank of the present invention.

[0046] Figure 14This describes the γ′ phase characteristics of the DP-HT-OC-9 small ingot blank of the present invention. Detailed Implementation

[0047] The technical solution of the present invention will be clearly and completely described below with reference to the accompanying drawings and specific embodiments. However, those skilled in the art will understand that the embodiments described below are only some embodiments of the present invention and are used only to illustrate the present invention, and should not be regarded as limiting the scope of the present invention. Where specific conditions are not specified in the embodiments, conventional conditions or conditions recommended by the manufacturer shall be followed. Where the manufacturers of reagents or instruments are not specified, they are all conventional products that can be purchased commercially.

[0048] In some embodiments of the present invention, a nickel-based superalloy is provided, comprising the following components by mass percentage:

[0049] C 0.05%–0.063%, Co 15.0%–15.04%, Cr 10.08%–10.13%, W 5.74%–5.81%, Mo 3.30%–3.34%, Al 3.73%–3.78%, Ti 3.68%–3.73%, Nb 1.77%–1.81%, Hf 0.20%–0.40%, B 0.010%–0.014%, Zr 0.008%–0.01%, Mg 0.002%–0.006%, balance Ni.

[0050] In different implementation methods, the mass percentages of each component in the nickel-based superalloy can be as follows:

[0051] The mass percentage of C can be 0.05%, 0.052%, 0.054%, 0.056%, 0.058%, 0.060%, 0.062%, or any value between two of these; the mass percentage of Co can be 15.0%, 15.01%, 15.02%, 15.03%, 15.04%, or any value between two of these; the mass percentage of Cr can be 10.08%, 10.09%, 10.10%, 10.11%, 10.12%, 10.13%, or any value between two of these; the mass percentage of W can be 5.74%, 5.7%, or any value between two of these. The mass percentage of Mo can be 5%, 5.76%, 5.77%, 5.78%, 5.79%, 5.80%, 5.81%, or any two of these values; the mass percentage of Mo can be 3.30%, 3.31%, 3.32%, 3.33%, 3.34%, or any two of these values; the mass percentage of Al can be 3.73%, 3.74%, 3.75%, 3.76%, 3.77%, 3.78%, or any two of these values; the mass percentage of Ti can be 3.68%, 3.69%, 3.70%, 3.71%, 3.72%, 3.73%, or... The mass percentage of Nb can be 1.77%, 1.78%, 1.79%, 1.80%, 1.81%, or any two of these values; the mass percentage of Hf can be 0.20%, 0.22%, 0.24%, 0.26%, 0.28%, 0.30%, 0.32%, 0.34%, 0.36%, 0.38%, 0.40%, or any two of these values, preferably 0.24% to 0.30%, more preferably 0.27% to 0.30%; the mass percentage of B can be 0.010%, 0.011%, 0.0%, or any two of these values. The mass percentage of Zr can be 0.008%, 0.0082%, 0.0084%, 0.0086%, 0.0088%, 0.009%, 0.0092%, 0.0094%, 0.0096%, 0.0098%, 0.01%, or any two of these values, preferably 0.0086% to 0.01%; the mass percentage of Mg can be 0.002%, 0.003%, 0.004%, 0.005%, 0.006%, or any two of these values.

[0052] In some embodiments, the mass ratio of Ti to Al in the nickel-based superalloy is (0.97 to 1):1; typically, but not limitingly, the mass ratio of Ti to Al can be a range of 0.97:1, 0.98:1, 0.99:1, 1:1, or any combination thereof.

[0053] In some embodiments, the total mass percentage of Al, Ti, Nb, and Hf in the nickel-based superalloy is 9.38% to 9.71%; typically, but not limitingly, for example, the total mass percentage of Al, Ti, Nb, and Hf in the nickel-based superalloy can be a range of 9.38%, 9.40%, 9.50%, 9.60%, 9.70%, 9.71%, or any combination thereof; preferably 9.47% to 9.54%.

[0054] The nickel-based superalloy of this invention is based on the composition of FGH4097 alloy. By increasing the content of C, Cr, and W and decreasing the content of Co and Mo, the solid solution strengthening effect is improved. By increasing the Ti content and decreasing the Al and Nb content, the mass ratio of Ti to Al is made close to 1:1. The total mass percentage of the main constituent elements of the γ′ phase, Al, Ti, Nb, and Hf, is 9.38% to 9.71%. Under the condition that the total amount of the γ′ phase forming elements Al, Ti, Nb, and Hf remains basically unchanged, the complete solid solution temperature of the γ′ phase is increased to 1188 to 1205°C, and the mass percentage of the γ′ phase is 57% to 59%.

[0055] The nickel-based superalloy of the present invention has better overall properties such as tensile properties, creep rupture properties, impact properties and fatigue properties than FGH4097 alloy, especially its tensile strength and high-temperature creep rupture properties are higher than those of FGH4097 alloy.

[0056] In some embodiments, the microstructure of the nickel-based superalloy is as follows:

[0057] The equilibrium phases in nickel-based superalloys include the matrix γ phase, γ' phase, and carbides (MC, M6C, M...). 23 C6), borides (M3B2, MB2) and trace amounts of TCP phase (μ phase, σ phase);

[0058] The average grain size of the nickel-based superalloy is 25–35 μm (e.g., 25 μm, 27 μm, 29 μm, 31 μm, 33 μm, 55 μm, etc.); the mass percentage of the γ′ phase in the nickel-based superalloy is 57%–59%;

[0059] The nickel-based superalloy has a primary γ′ phase with an average size of 0.5–4 μm (e.g., 0.5 μm, 1 μm, 1.5 μm, 2 μm, 2.5 μm, 3 μm, 3.5 μm, 4 μm, etc.), a secondary γ′ phase with an average size of 0.2–0.8 μm (e.g., 0.2 μm, 0.4 μm, 0.6 μm, 0.8 μm, etc.), and a tertiary γ′ phase with an average size of 0.04–0.07 μm (e.g., 0.04 μm, 0.05 μm, 0.06 μm, 0.07 μm, etc.). In some embodiments, the nickel-based superalloy exhibits the following properties:

[0060] Room temperature (25°C) tensile properties of nickel-based superalloys: Room temperature tensile strength is 1510–1590 MPa (e.g., 1510 MPa, 1520 MPa, 1530 MPa, 1540 MPa, 1550 MPa, 1560 MPa, 1570 MPa, 1580 MPa, 1590 MPa, etc.), and room temperature yield strength is 1020–1185 MPa (e.g., 1020 MPa, 104 MPa, etc.). 0MPa, 1060MPa, 1080MPa, 1100MPa, 1120MPa, 1140MPa, 1160MPa, 1180MPa, 1185MPa), room temperature elongation at break is 14% to 25% (e.g., 14%, 16%, 18%, 20%, 22%, 25%), room temperature reduction of area is 17% to 23% (e.g., 17%, 19%, 21%, 23%, etc.);

[0061] High-temperature (750℃) tensile properties of nickel-based superalloys: tensile strength of 1090–1250 MPa (e.g., 1090 MPa, 1110 MPa, 1130 MPa, 1150 MPa, 1170 MPa, 1190 MPa, 1210 MPa, 1230 MPa, 1250 MPa, etc.), yield strength of 913–1050 MPa (e.g., 913 MPa, 1000 MPa, 1020 MPa, 1040 MPa, 1050 MPa, etc.), elongation after fracture of 12%–35% (e.g., 12%, 15%, 20%, 25%, 30%, 35%, etc.), reduction of area of ​​12%–32% (e.g., 12%, 15%, 20%, 25%, 30%, 32%, etc.);

[0062] Room temperature impact properties of nickel-based superalloys: energy absorption of 27–42 J (e.g., 27 J, 29 J, 31 J, 33 J, 35 J, 37 J, 39 J, 41 J, 42 J, etc.), and impact toughness of 33–53 J / cm. 2 (For example, 33J / cm) 2 35J / cm 2 37J / cm 239J / cm 2 41J / cm 2 43J / cm 2 45J / cm 2 47J / cm 2 49J / cm 2 51J / cm 2 53J / cm 2 etc);

[0063] Low-cycle fatigue properties of nickel-based superalloys: fatigue life of 11,000 to 50,000 N (11,000 N, 20,000 N, 30,000 N, 40,000 N, 50,000 N, etc.) at 650℃ and 1100 MPa.

[0064] Smooth creep properties of nickel-based superalloys at 650℃ and 1140MPa: fracture time is 7 to 61 hours (e.g., 7 hours, 10 hours, 20 hours, 30 hours, 40 hours, 50 hours, 60 hours, etc.), elongation after fracture is 8% to 27% (e.g., 8%, 10%, 14%, 18%, 22%, 25%, 27%, etc.), and reduction of area is 8% to 24% (e.g., 8%, 10%, 14%, 18%, 22%, 24%, etc.).

[0065] Smooth creep properties of nickel-based superalloys at 750℃ and 750MPa: fracture time is 29–130 h (e.g., 29 h, 50 h, 70 h, 90 h, 110 h, 130 h, etc.), elongation after fracture is 4.5%–15% (e.g., 4.5%, 6%, 8%, 10%, 12%, 14%, 15%, etc.), and reduction of area is 5%–18% (e.g., 5%, 7%, 9%, 11%, 13%, 15%, 17%, 18%, etc.).

[0066] Composite creep rupture properties of nickel-based superalloys: The composite creep rupture time at 650℃ and 990MPa is 180–220 h (e.g., 180 h, 200 h, 220 h, etc.), and the composite creep rupture time at 750℃ and 637MPa is 190–210 h (e.g., 190 h, 200 h, 210 h, etc.).

[0067] In some embodiments of the present invention, a method for preparing the above-mentioned nickel-based high-temperature alloy is also provided, comprising the following steps:

[0068] Alloy powder is prepared by mixing ingredients according to the specified ratio; the alloy powder is then subjected to hot isostatic pressing, solution treatment and aging treatment in sequence to obtain a nickel-based high-temperature alloy.

[0069] In some embodiments, the alloy powder is prepared by: preparing master alloy bars by vacuum horizontal continuous casting (VHCC) after batching according to the specified ratio; and then atomizing the master alloy bars into alloy powder by plasma rotating electrode method (PREP).

[0070] In some embodiments, the particle size of the alloy powder is ≤100μm; for example, the particle size of the alloy powder is 50-100μm or the particle size of the alloy powder is ≤75μm.

[0071] In some embodiments, the alloy powder is sequentially sieved, electrostatically separated, and inspected before being packed into a sleeve for degassing, sleeve assembly, and sealing, followed by hot isostatic pressing.

[0072] In some embodiments, the hot isostatic pressing (HIP) treatment includes holding at 1190–1220°C and 130–150 MPa for 3–5 hours. Typically, but not limitingly, the HIP temperature can be a range of 1190°C, 1200°C, 1210°C, 1220°C, or any combination thereof. Preferably, the HIP temperature is 1200°C and the HIP time is 4 hours. Using the above-described HIP treatment, this invention can obtain nickel-based superalloys with better γ′ phase distribution, quantity, size, and morphology.

[0073] In some embodiments, the solution treatment includes: holding at 1170–1220°C for 2–5 hours, followed by cooling; typically, but not limitingly, for example, the solution treatment temperature can be a range of 1170°C, 1180°C, 1190°C, 1200°C, 1210°C, 1220°C, or any combination thereof; preferably, the solution treatment temperature is 1200°C, and the holding time is 4 hours.

[0074] In some embodiments, the cooling method for the solution treatment includes furnace cooling followed by air cooling; or, the cooling method for the solution treatment includes oil cooling.

[0075] In some embodiments, the cooling rate of the solution treatment is 30–40 °C / min; or, the cooling rate of the solution treatment is ≥80 °C / min; preferably 80–110 °C / min; typically, but not limitingly, for example, the cooling rate of the solution treatment can be 30 °C / min, 35 °C / min, or 40 °C / min; or, it can be a range of 85 °C / min, 90 °C / min, 100 °C / min, 110 °C / min, or any combination thereof.

[0076] The cooling method and cooling rate of solution treatment have a significant impact on the grain size and properties of the obtained nickel-based superalloys. By adopting the above-mentioned cooling method and cooling rate, the high-temperature creep performance of the alloy can be further improved, the creep life can be increased, and the problem of composite creep notch sensitivity can be solved.

[0077] In some embodiments, the aging treatment includes: heat treatment at 870±10℃ for 6–9 hours, followed by air cooling, and then heat treatment at 760±10℃ for 14–18 hours, followed by air cooling. Preferably, the aging treatment of the present invention is a two-stage aging treatment of heat treatment at 870℃ for 8 hours, heat treatment at 760℃ for 16 hours, followed by air cooling.

[0078] The microstructure and properties of nickel-based superalloys obtained by sequentially subjecting alloy powders with a particle size of 50–100 μm to hot isostatic pressing at 1200℃, solution treatment (temperature 1180–1210℃, cooling method: furnace cooling followed by air cooling), and aging treatment are as follows:

[0079] The average grain size is 28.32–34.95 μm; the average size of the primary γ′ phase is 2.58–2.63 μm, the average size of the secondary γ′ phase is 0.61–0.70 μm, and the average size of the tertiary γ′ phase is 0.06–0.07 μm.

[0080] Room temperature (25℃) tensile properties of nickel-based superalloys: tensile strength 1514~1557MPa, yield strength 1024~1068MPa, elongation after fracture 19.5%~24.5%, reduction of area 20%~23%;

[0081] High-temperature (750℃) tensile properties of nickel-based superalloys: tensile strength is 1090~1160MPa, yield strength is 913~1000MPa, elongation after fracture is 23%~35%, and reduction of area is 17%~32%.

[0082] Composite creep rupture properties of nickel-based superalloys: The composite creep rupture time is 210-220 h at 650℃ and 990 MPa, and 190-200 h at 750℃ and 637 MPa.

[0083] The microstructure and properties of nickel-based superalloys obtained by sequentially subjecting alloy powder with a particle size of 50–100 μm to hot isostatic pressing at 1200℃, solution treatment (temperature 1180–1210℃, cooling method: oil cooling), and aging treatment are as follows:

[0084] The average grain size is 32–33 μm; the average size of the primary γ′ phase is 0.60–0.70 μm, the average size of the secondary γ′ phase is 0.20–0.30 μm, and the average size of the tertiary γ′ phase is 0.06–0.07 μm.

[0085] Room temperature (25℃) tensile properties of nickel-based superalloys: tensile strength 1556~1587MPa, yield strength 1129~1185MPa, elongation after fracture 14.5%~17.5%, reduction of area 17~18%;

[0086] High-temperature (750℃) tensile properties of nickel-based superalloys: tensile strength is 1200~1250MPa, yield strength is 990~1050MPa, elongation after fracture is 12%~24%, and reduction of area is 13%~25.5%.

[0087] Composite creep rupture properties of nickel-based superalloys: The composite creep rupture time is 180-190 h at 650℃ and 990 MPa, and 200-210 h at 750℃ and 637 MPa.

[0088] The microstructure and properties of nickel-based superalloys obtained by sequentially subjecting alloy powder with a particle size ≤75μm to hot isostatic pressing at 1200℃, solution treatment (temperature 1180~1210℃, cooling method: oil cooling) and aging treatment are as follows:

[0089] The average grain size is 25–26 μm;

[0090] The average size of the primary γ′ phase is 0.6–0.7 μm, the average size of the secondary γ′ phase is 0.2–0.3 μm, and the average size of the tertiary γ′ phase is 0.045–0.055 μm.

[0091] Room temperature (25℃) tensile properties of nickel-based superalloys: tensile strength 1563~1568MPa, yield strength 1138~1141MPa, elongation after fracture 18.5%~19.5%, reduction of area 19%~21%;

[0092] The high-temperature (750℃) tensile properties of the nickel-based superalloy are as follows: tensile strength of 1220–1250 MPa, yield strength of 1030–1050 MPa, elongation after fracture of 12%–19%, and reduction of area of ​​12%–17%. The preparation method of the nickel-based superalloy of this invention, through research on hot isostatic pressing and heat treatment processes, improves the performance of the nickel-based superalloy. The room-temperature tensile strength, high-temperature tensile strength, and creep strength of the obtained nickel-based superalloy are all superior to those of the FGH4097 alloy prepared by direct hot isostatic pressing near-net-shape forming. It can meet the requirements of aero-engines for 750℃ disc materials and is applicable to high-end components of different types of aero-engines and ground gas turbines, such as discs, shafts, and ring components.

[0093] The low-cost, high-performance, high-reliability, high-corrosion-resistant, and long-life nickel-based superalloys prepared by this invention can be applied to higher-generation new carrier-based aircraft, carrier-based anti-submarine aircraft, carrier-based early warning aircraft, carrier-based electronic warfare aircraft, reconnaissance aircraft, tanker aircraft, transport aircraft, early warning / reconnaissance helicopters, anti-ship helicopters, mine countermeasures helicopters, search and rescue helicopters, etc., as well as gas turbines for ground power generation or other industrial fields, including aero-engine gas turbines and aero-derivative gas turbines.

[0094] Examples 1-4

[0095] The method for preparing nickel-based superalloys provided in this embodiment yields nickel-based superalloys with the following composition (wt%) as shown in the table.

[0096] Table 1

[0097] C Co Cr W Mo Al Ti Nb Hf B Zr Mg Ni Example 1 0.058 15.04 10.10 5.80 3.30 3.73 3.68 1.8 0.26 0.010 0.01 0.002 Balance Example 2 0.060 15.02 10.08 5.75 3.34 3.77 3.72 1.78 0.24 0.012 0.009 0.004 Balance Example 3 0.062 15.0 10.12 5.77 3.32 3.75 3.70 1.79 0.30 0.014 0.008 0.006 Balance Example 4 0.056 15.01 10.09 5.79 3.31 3.78 3.71 1.77 0.28 0.013 0.009 0.005 Balance

[0098] Specifically, the preparation method of nickel-based superalloys includes the following steps:

[0099] After the materials are proportioned, the master alloy bar is prepared by vacuum horizontal continuous casting (VHCC); the master alloy bar is atomized and powdered by plasma rotating electrode method (PREP) to obtain alloy powder with a particle size of 50-100μm; the alloy powder is subjected to hot isostatic pressing treatment, then the cladding is removed, and it is divided into four parts in a longitudinal cross shape. Each part is then cut into 1 / 4 pieces and subjected to solution treatment and aging treatment in sequence to obtain small ingots.

[0100] Among them, the hot isostatic pressing treatment is: heat treatment at 1200℃ and 140MPa for 4 hours;

[0101] Solution treatment was performed by holding at 1180–1210°C for 4 hours, followed by cooling. The cooling method was furnace cooling followed by air cooling. The cooling rate was 30–40°C / min. The solution treatment temperatures for Examples 1–4 were 1180°C, 1190°C, 1200°C, and 1210°C, respectively. The aging treatment included holding at 870°C for 8 hours, followed by air cooling, then holding at 760°C for 16 hours, followed by air cooling.

[0102] Comparative Examples 1-2

[0103] The preparation process of the nickel-based superalloy provided in this comparative example yields nickel-based superalloys with the following composition (wt%) as shown in Table 2.

[0104] Table 2

[0105] C Co Cr W Mo Al Ti Nb Hf B Zr Mg Ce Ta Ni Comparative Example 1 0.05 16.0 9.5 5.5 3.85 5.0 1.80 2.6 0.3 0.01 0.013 0.02 0.009 - Balance Comparative Example 2 0.05 8.0 13.0 3.6 3.5 3.5 2.5 3.6 - 0.01 0.05 - - 0.2 Balance

[0106] The preparation method of the above-mentioned nickel-based superalloy is described in Example 1.

[0107] Experimental Example 1

[0108] The grain structure of the small ingots prepared in Examples 1, 2, 3, and 4 is as follows: Figure 1 , Figure 2 , Figure 3 and Figure 4 As shown; Figures 1 to 4 In Figure a, the scale bar is 100 μm, and in Figure b, the scale bar is 20 μm. The γ′ phase characteristics of the small ingots prepared in Examples 1, 2, 3, and 4 at different magnifications are shown below. Figure 5 , Figure 6 , Figure 7 and Figure 8 As shown.

[0109] The grain size and γ′ phase characteristics of the small ingots prepared in Examples 1, 2, 3 and 4 are shown in Table 3.

[0110] Table 3

[0111]

[0112] from Figures 1 to 4 As shown in Table 3, the nickel-based superalloys in Examples 1-4 exhibit a gradual increase in grain size after the same aging heat treatment with increasing solution temperature (1180-1210℃), with an average grain size of 28.32-34.95 μm and a grain size grade of 6.0-7.0. At a solution temperature of 1180℃, recrystallization is insufficient, resulting in an uneven grain size distribution with a bimodal distribution, and the presence of a small amount of PPB and some residual dendrites. The nickel-based superalloys in Comparative Examples 1-2 all have a grain size grade of 6.0-6.5.

[0113] from Figures 5 to 8As shown in Table 3, under the same solution treatment holding time, cooling method and two-stage aging, the nickel-based superalloys in Examples 1 to 4 showed little change in the size of the primary γ′ phase after aging treatment as the solution treatment temperature increased (1180 to 1210 °C), with an average size of 2.58 to 2.63 μm, mainly in the form of irregular elongated strips, distributed on the grain boundaries. At a solution treatment temperature of 1180℃, the number of primary γ′ phases at the grain boundaries is relatively large. As the solution treatment temperature increases, the number of primary γ′ phases gradually decreases, while the size of secondary γ′ phases gradually increases, with an average size of 0.61–0.70 μm. At 1180℃, the secondary γ′ phases are mainly triangular and butterfly-shaped, while at 1190–1210℃, they are mainly cubic and grid-shaped, and are undergoing splitting, distributed within the grains. The tertiary γ′ phases have similar sizes, with an average size of 0.07 μm, and are all small spherical, mainly distributed between the primary and secondary γ′ phases. In summary, the small ingot treated at 1200℃ for 4 hours has better results. In Comparative Example 1, the γ′ phase size in the nickel-based superalloy is smaller than that in Examples 1–4, while in Comparative Example 2, the γ′ phase size is close to or slightly larger than that in Examples 1–4.

[0114] The properties of the small ingots prepared in Examples 1-4 and Comparative Examples 1-2 were tested, and the results are shown in Tables 4, 5, 6 and 7.

[0115] Table 4

[0116]

[0117] As shown in Table 4, the room temperature tensile properties of the nickel-based superalloys in Examples 1-4, when the solution temperature is 1180℃ (Example 1), exhibit a relatively high average room temperature tensile strength of 1547 MPa, but a slightly lower average yield strength of 1028 MPa. Comparing the overall strength and plasticity, when the solution temperature is 1200-1210℃ (Examples 2-4), the relatively high average room temperature tensile strengths are 1532 MPa and 1532 MPa, respectively; the relatively high average yield strengths are 1062 MPa and 1030 MPa, respectively; the average elongation after fracture is 24% and 24%, respectively; and the average reduction of area is 22% and 23%, respectively. Compared with Comparative Example 1, the nickel-based superalloys in Examples 1-4 exhibit higher room temperature tensile strength and plasticity; compared with Comparative Example 2, their strength is slightly lower, but their plasticity is higher.

[0118] As shown in Table 4, the high-temperature tensile properties of the nickel-based superalloys in Examples 1-4 increase with increasing solution temperature (1180-1210℃), resulting in higher strength and ductility. At solution temperatures of 1200-1210℃, the highest average tensile strengths were 1113 MPa and 1145 MPa, respectively; the highest average yield strengths were 924 MPa and 1000 MPa, respectively; the average elongation after fracture was 34% and 32%, respectively; and the average reduction of area was 29% and 32%, respectively. Compared to Comparative Example 1, the nickel-based superalloys in Examples 3-4 exhibited similar tensile strength and ductility at 750℃.

[0119] Table 5

[0120]

[0121] As shown in Table 5, the nickel-based superalloys in Examples 1-4 exhibited the lowest impact energy and impact toughness at a solution treatment temperature of 1180℃, with average values ​​of 35 J and 43 J / cm, respectively. 2 When the solution treatment temperature is between 1190 and 1210℃, the alloy exhibits high impact energy and impact toughness, with average values ​​ranging from 38 to 42 J and 48 to 52 J / cm, respectively. 2 Compared with Comparative Examples 1 and 2, the nickel-based superalloys in Examples 1 to 4 have higher room temperature impact energy and impact toughness.

[0122] As shown in Table 5, the low-cycle fatigue properties of the nickel-based superalloys in Examples 1-4 exhibit low-cycle fatigue lives ranging from 11,119 to 33,012 cycles at 650℃ / 1100MPa. The longest low-cycle fatigue life (33,012 cycles) is achieved when the solution temperature is 1200℃. Comparative Examples 1-2 show low-cycle fatigue lives of 63,000 and 16,400 cycles under different conditions: 650℃ / 980MPa and 540℃ / strain ratio 0.95, respectively.

[0123] Table 6

[0124]

[0125]

[0126] As shown in Table 6, the creep rupture properties at 650℃ / 1140MPa indicate that the creep rupture life and plasticity of the nickel-based superalloys in Examples 1-4 do not change significantly with increasing solution temperature (1180-1210℃). Among them, the creep rupture life is relatively high at 9.74h when the solution temperature is 1200℃, with elongation after fracture and reduction of area of ​​25% and 23% respectively, which reduces stress. The creep rupture life of Comparative Examples 1-2 at 650℃ / 1120MPa and 650℃ / 1035MPa are 10h and 80h respectively, and both have low plasticity.

[0127] As can be seen from Table 6 on the smooth creep properties at 750℃ / 750MPa, the creep life of the nickel-based superalloys in Examples 1 to 4 increases with the increase of solution temperature (1180 to 1210℃). Among them, when the solution temperature is 1200 to 1210℃, the creep life of the alloy is relatively high, at 63.6h and 69.3h respectively, and the plasticity change is not obvious.

[0128] Table 7

[0129]

[0130] As can be seen from Table 7 on the composite creep rupture performance at 650℃ / 1140MPa, the creep rupture life of the nickel-based superalloy in Example 3 is 98h longer than that of Comparative Example 1, and the plasticity is comparable. As can be seen from Table 7 on the composite creep rupture performance at 750℃ / 750MPa, the creep rupture life of the nickel-based superalloy in Example 3 is 65h longer than that of Comparative Example 1, and the plasticity is comparable.

[0131] Experimental Example 2

[0132] The effects of hot isostatic pressing (hot isostatic pressing temperature) and heat treatment processes (solution treatment cooling rate: furnace cooling + air cooling (cooling rate of 30-40℃ / min), oil cooling (cooling rate ≥80℃ / min)) on the grain size, quantity, size, morphology and distribution characteristics of alloy powders with different particle sizes (50-100μm and ≤75μm) were investigated.

[0133] Master alloy bars were prepared using vacuum horizontal continuous casting (VHCC) technology, and nickel-based superalloy powders of different particle sizes were prepared using plasma rotating electrode method (PREP). The composition of the nickel-based superalloy powders was the same as that in Example 3. After being degassed, packed, and sealed in a cladding, the grain size and grain level statistical results after hot isostatic pressing at different temperatures of 1210℃ and 1200℃ and 130MPa for 4 hours are shown in Table 8, and the γ′ phase characteristics are shown in Table 9.

[0134] Table 8

[0135]

[0136] Table 8 shows that after hot isostatic pressing at 1200℃ and 1210℃, the alloy powder with a particle size of 50–100 μm has an average grain size of 31.74–32.53 μm and a grain size of 6.5–7.0, respectively. The grain size distribution is relatively uniform, and recrystallization is relatively complete, with a small amount of original particle boundaries (PPB) and residual dendrites present. After hot isostatic pressing at 1200℃ and 1210℃, the alloy powder with a particle size ≤75 μm has an average grain size of 28.42–29.49 μm and a grain size of 6.5–7.0, respectively. The grain size is finer and more uniform than that of the alloy powder with a particle size of 50–100 μm after hot isostatic pressing, and recrystallization is relatively complete, with a small amount of PPB and residual dendrites present.

[0137] Table 9

[0138]

[0139] As shown in Table 9, after hot isostatic pressing (HIP) at 1200℃ and 1210℃, the size of the primary γ′ phase of alloy powder with a particle size of 50-100μm did not change much. The primary γ′ phase was characterized by irregular elongated shape. However, the number of primary γ′ phases after HIP at 1200℃ was greater than that after HIP at 1210℃, and they were mainly distributed on the grain boundaries. The size of the secondary γ′ phase was slightly smaller than that after HIP at 1210℃. The secondary γ′ phase was characterized by a large number of precipitates, mainly cubic and butterfly-shaped, and was undergoing splitting. The size of the tertiary γ′ phase was slightly larger than that after HIP at 1210℃. The tertiary γ′ phase was characterized by fine spherical shape, and there were two types of tertiary γ′ phases with different sizes: large and small. After hot isostatic pressing (HIP) at 1200℃ and 1210℃, the sizes of the primary, secondary, and tertiary γ′ phases of alloy powders with a particle size ≤75μm showed little change. The primary γ′ phases were all irregularly elongated, but the number of primary γ′ phases was greater after HIP at 1200℃ than at 1210℃, and they were mainly distributed at grain boundaries. The secondary γ′ phase precipitated in greater quantities, mainly in cubic and butterfly shapes, and was undergoing splitting. The tertiary γ′ phase was finely spherical, with both large and small sizes. After HIP, the primary γ′ phase of alloy powders with a particle size ≤75μm was slightly smaller than that of alloy powders with a particle size of 50–100μm, the secondary γ′ phase was slightly larger, and the tertiary γ′ phase was roughly the same size. Alloy powders with particle sizes of 50–100 μm and ≤75 μm showed good distribution, quantity, size, and morphology of the three γ′ phases after hot isostatic pressing at 1200 °C.

[0140] After removing the outer casing of BT-3, the small ingot blank is longitudinally divided into four parts in a cross shape. Each part is then cut into quarter pieces and subjected to solution treatment and secondary aging treatment (held at 870℃ for 8 hours, air-cooled, held at 760℃ for 16 hours, and air-cooled). The resulting ingot blank is named DP-HT-3, with a diameter and height of Φ80~90mm×h130~140mm.

[0141] After removing the outer casing, small ingots of BT-1, BT-3, and BT-9 were longitudinally divided into four quarters in a cross shape. Each quarter was then subjected to solution treatment (oil cooling) and secondary aging treatment (held at 870℃ for 8 hours, air-cooled, then held at 760℃ for 16 hours, and air-cooled). The resulting ingots were named DP-HT-OC-1, DP-HT-OC-3, and DP-HT-OC-9, respectively. Their diameter and height dimensions were Φ80~90mm×h130~140mm.

[0142] The grain size and grain level of each of the above-mentioned small ingots are recorded in Table 12, and the γ′ phase characteristics of each small ingot are recorded in Table 13; the grain structure of the DP-HT-3 small ingot is as follows: Figure 3 As shown, its γ′ phase characteristics are as follows Figure 7 As shown, the same as in Example 3; the grain structures of DP-HT-OC-1, DP-HT-OC-3, and DP-HT-OC-9 small ingots are as follows. Figure 9 , Figure 10 , Figure 11 As shown ( Figures 9 to 11 In Figure a, the scale bar is 100 μm, and in Figure b, the scale bar is 20 μm. The γ′ phase characteristics of DP-HT-OC-1, DP-HT-OC-3, and DP-HT-OC-9 ingots at different magnifications are shown below. Figure 12 , Figure 13 , Figure 14 As shown.

[0143] Table 12

[0144]

[0145] from Figures 9 to 11 As shown in Table 12, the grain size of DP-HT-OC-1, DP-HT-OC-3, and DP-HT-OC-9 small ingots after solution treatment and oil cooling is smaller than that after solution treatment and furnace cooling + air cooling. The grain size of DP-HT-OC-1 small ingot after oil cooling is approximately 32.69 μm, with a grain size of 6.5–7.0 grade. After furnace cooling + air cooling, the grain size is slightly larger, approximately 34.25 μm. The grain size of DP-HT-OC-3 small ingot after oil cooling is approximately 32.92 μm, with a grain size of 6.5–7.0 grade. After furnace cooling + air cooling, the grain size is approximately 33.67 μm, with a grain size of 6.0–6.5 grade. The grain size of the DP-HT-OC-9 small ingot after oil cooling is approximately 25.62 μm, with a grain size grade of 7.0–7.5. After furnace cooling followed by air cooling, the grain size is approximately 29.15 μm, with a grain size grade of 6.5–7.0. This demonstrates that the solution cooling rate has a significant impact on the grain size of the small ingot, and oil cooling can reduce the alloy grain size.

[0146] Table 13

[0147]

[0148] from Figures 12 to 14 As shown in Table 13, the γ′ phase size of DP-HT-OC-1, DP-HT-OC-3, and DP-HT-OC-9 small ingots treated with oil cooling is smaller than that treated with furnace cooling + air cooling. Taking DP-HT-OC-3 small ingots treated with oil cooling and DP-HT-OC-3 treated with furnace cooling + air cooling as examples for comparison, overall, the primary and secondary γ′ phase sizes of DP-HT-OC-3 small ingots treated with oil cooling are smaller than those treated with furnace cooling + air cooling, while the tertiary γ′ phase is close to or slightly smaller. After oil cooling, the primary γ′ phase in the DP-HT-OC-3 small ingot is smaller and irregularly elongated, mainly distributed on the grain boundaries. The secondary γ′ phase precipitates in greater quantities and is smaller in size, mainly distributed in cubic shapes within the grains. The tertiary γ′ phase precipitates in fewer quantities, appearing as small spheres, mainly distributed around the primary γ′ phase, and its size is close to or slightly smaller than that of furnace cooling + air cooling. The DP-HT-OC-1 and DP-HT-OC-9 small ingots were oil-cooled, and the γ′ phase characteristics were similar to those of the DP-HT-OC-3 small ingots that were oil-cooled.

[0149] FGH4097 Standard A (alloy composition can be found in the "Chinese High Temperature Alloy Standard Handbook (Volume 2)") heat treatment regime: solution heat treatment is (1120~1160)±10℃ / (2~4)h / salt quenching or oil quenching, first-stage aging is 870℃±10℃ / 1h / AC, and second-stage aging is 650℃±10℃ / 24h / AC; FGH4097 Standard B (alloy composition is the same as Standard A) heat treatment regime: solution heat treatment is (1140~1160)±10℃ / (1~2)h / salt quenching, first-stage aging is 870℃±10℃ / 1.5h / AC, and second-stage aging is 650℃±10℃ / 24h / AC.

[0150] The room temperature (25℃) tensile properties of each small ingot were tested, and the results are shown in Table 14. The high temperature (750℃) tensile properties of each small ingot were tested, and the results are shown in Table 15.

[0151] Table 14

[0152]

[0153] Table 14 shows that, comparing the room temperature tensile strength of alloys treated with the same solution treatment temperature of 1200℃, the same holding time of 4h, and different solution cooling rates (furnace cooling + air cooling, oil cooling) + aging treatment, the alloy treated with solution cooling and oil cooling has an average room temperature tensile strength 40MPa higher than that treated with furnace cooling + air cooling. The average yield strength of the former is 127MPa higher than the latter, while the plasticity of the former is slightly lower. The average elongation after fracture is 8% lower, and the average reduction of area is 6% lower. Comparing the room temperature tensile strength of alloys treated with solution cooling and oil cooling for different powder particle sizes (50–100μm and ≤75μm), both have similar strength and plasticity, with no significant difference. Overall, comparing the room temperature tensile properties of the alloy of this invention with FGH4097 alloy (two aging regimes), the room temperature tensile strength and plasticity of the alloy of this invention are higher than those of FGH4097 alloy.

[0154] Table 15

[0155]

[0156]

[0157] Table 15 shows the high-temperature tensile properties of the alloys. Comparing the tensile properties at 750℃ of the alloys after solution cooling at the same solution temperature of 1200℃, the same holding time of 4h, and different solution cooling rates (furnace cooling + air cooling, oil cooling) + aging treatment, the alloy treated with solution cooling and oil cooling has an average tensile strength 112MPa higher than that treated with furnace cooling + air cooling. The average yield strength of the former is 97MPa higher than that of the latter. The former has lower plasticity, the elongation after fracture is 16% lower than that of the latter, and the reduction of area is 10% lower than that of the latter. Comparing the high-temperature tensile properties at 750℃ of alloys with different powder particle sizes (50~100μm and ≤75μm) after solution cooling and oil cooling, the strength and plasticity of the two are similar, with no significant difference. Overall, comparing the high-temperature tensile properties at 750℃ of the alloy of this invention and the FGH4097 alloy (two aging conditions), the high-temperature tensile strength of the alloy of this invention treated with oil cooling is similar to that of the FGH4097 alloy, but the plasticity of the alloy of this invention treated with oil cooling is lower than that of FGH4097. The alloy produced by furnace cooling and air cooling has a lower high-temperature tensile strength than FGH4097 alloy, but its plasticity is slightly higher or similar to that of FGH4097 alloy.

[0158] The room temperature (25℃) impact properties of each small ingot were tested, and the results are shown in Table 16.

[0159] Table 16

[0160]

[0161] Table 16 shows that, comparing the impact energy and impact toughness of alloys treated at the same solution temperature of 1200℃, the same holding time of 4h, and different solution cooling rates (furnace cooling + air cooling, oil cooling) + aging treatment, the average impact energy and impact toughness of the alloy treated with solution cooling and oil cooling are 7J and 8J / cm lower, respectively, than those treated with furnace cooling + air cooling. 2 Comparing the impact energy and impact toughness of the alloys after solution treatment and oil cooling with different powder particle sizes (50–100 μm and ≤75 μm), the impact energy and impact toughness of the two were similar, with no significant difference. Overall, comparing the room temperature impact properties of the alloy of this invention with FGH4097 alloy (two aging regimes), the room temperature impact energy and impact toughness of the oil-cooled alloy of this invention were slightly lower than those of FGH4097 alloy. The room temperature impact energy and impact toughness of the alloy of this invention, which used furnace cooling + air cooling, were similar to those of FGH4097 alloy.

[0162] The low-cycle fatigue properties of each small ingot were tested, and the results are shown in Table 17.

[0163] Table 17

[0164]

[0165] Table 17 shows that, comparing the low-cycle fatigue performance of alloys treated with the same solution treatment temperature of 1200℃, the same holding time of 4h, and different solution cooling rates (furnace cooling + air cooling, oil cooling) + aging treatment, the alloy treated with solution cooling and oil cooling has a low-cycle fatigue life 6791 cycles lower than that treated with furnace cooling + air cooling. Comparing the low-cycle fatigue performance of alloys with different powder particle sizes (50–100μm and ≤75μm) treated with solution cooling and oil cooling, the alloy with ≤75μm powder particle size has a low-cycle fatigue life 14516 cycles higher than that with 50–100μm powder particle size. Overall, comparing the low-cycle fatigue performance of the alloy of this invention with that of FGH4097 alloy (two aging conditions), the low-cycle fatigue life of the alloy of this invention after oil cooling or furnace cooling + air cooling is slightly higher than that of the FGH4097 alloy.

[0166] The smoothness and durability of each small ingot were tested, and the results are shown in Tables 18 and 19.

[0167] Table 18

[0168]

[0169]

[0170] Table 18 shows that, comparing the smooth creep rupture properties of alloys treated with the same solution temperature of 1200℃, the same holding time of 4h, and different solution cooling rates (furnace + air cooling and oil cooling) + aging treatment at 650℃ / 1140MPa, the alloy treated with solution cooling and oil cooling has a creep rupture life 27.86h longer than that treated with furnace cooling + air cooling. The former has lower plasticity than the latter, with elongation after fracture and reduction of area being 18.5% and 14.8% lower, respectively. Comparing the smooth creep rupture properties of alloys with different powder particle sizes (50–100μm and ≤75μm) after solution cooling and oil cooling, the alloy with 50–100μm powder particle size has a creep rupture life 19h longer than that with ≤75μm powder particle size. However, the former's coarser powder has worse plasticity than the latter's finer powder, with elongation after fracture and reduction of area being 5.5% and 3.8% lower, respectively. After reducing the creep rupture stress, the creep rupture performance at 650℃ / 1100MPa was compared with that of alloys treated at the same solution temperature of 1200℃, the same holding time of 4h, and different solution cooling rates (furnace cooling + air cooling, oil cooling) + aging treatment at 650℃ / 1100MPa. The creep rupture life of the alloys after stress reduction was extended. The creep rupture life of the alloy treated with solution cooling and oil cooling was 61h, which is 42.9h longer than that treated with furnace cooling + air cooling. The former had lower plasticity than the latter, with elongation after fracture and reduction of area being 12.7% and 13.2% lower, respectively. Comparing the creep rupture performance of the alloy of this invention with that of FGH4097 alloy (first-order aging regime) at 650℃ / 1025~1120MPa after stress reduction, the creep rupture life of the alloy of this invention after oil cooling or furnace cooling + air cooling was higher than that of the FGH4097 alloy.

[0171] Table 19

[0172]

[0173]

[0174] Table 19 shows that, comparing the smooth creep rupture properties of alloys treated with the same solution temperature of 1200℃, the same holding time of 4h, and different solution cooling rates (furnace + air cooling and oil cooling) + aging treatment at 750℃ / 750MPa, the alloy treated with solution cooling and oil cooling has a creep rupture life 12.9h shorter than that treated with furnace cooling + air cooling. The former also has lower plasticity, and its elongation after fracture and reduction of area are 7.7% and 8.9% lower, respectively. Comparing the smooth creep rupture properties of alloys with different powder particle sizes (50–100μm and ≤75μm) after solution cooling and oil cooling, the creep rupture life and plasticity are similar for both. After reducing the creep rupture stress, the creep rupture performance at 750℃ / 690MPa was compared with that of the alloys after the same solution temperature of 1200℃, the same holding time of 4h, and different solution cooling rates (furnace cooling + air cooling, oil cooling) + aging treatment at 750℃ / 690MPa. The creep rupture life of the alloys after stress reduction was extended in all cases. The creep rupture life of the oil-cooled alloy was 130h, similar to that of the furnace-cooled + air-cooled alloy, and their plasticity was also similar. Comparing the creep rupture performance of the alloy of this invention with that of FGH4097 alloy (first-order aging regime) at 750℃ / 690MPa after stress reduction, the creep rupture life of the alloy of this invention after oil cooling or furnace cooling + air cooling was higher than that of the FGH4097 alloy.

[0175] The composite durability properties of each small ingot were tested, and the results are shown in Tables 20 and 21.

[0176] Table 20

[0177]

[0178] Table 20 shows that the creep rupture life and plasticity of the oil-cooled alloy are 32 hours lower than those of the furnace-cooled + air-cooled alloy. Its elongation after fracture and reduction of area are also 4.9% and 4.8% lower, respectively, compared to the furnace-cooled + air-cooled alloy. Comparing the composite creep rupture performance of the alloy of this invention with that of FGH4097 alloy (two aging regimes) at 650℃ / 980~990MPa, the composite creep rupture life of the alloy of this invention after oil cooling or furnace-cooled + air cooling is higher than that of the FGH4097 alloy.

[0179] Table 21

[0180]

[0181]

[0182] Table 21 shows that, comparing the composite creep rupture properties of the alloys at 750℃ / 637MPa after the same solution treatment temperature (1200℃), the same holding time (4h), and different solution cooling rates (furnace cooling + air cooling, oil cooling) + aging treatment, the creep rupture life of the alloy treated with oil cooling is 203.5h, which is similar to that of the alloy treated with furnace cooling + air cooling. However, if the cooling rate of oil cooling is too fast, the alloy is prone to fracture notches. Comparing the composite creep rupture properties of the alloy of this invention and the FGH4097 alloy (first-order aging regime) at 750℃ / 637MPa, the composite creep rupture life of the alloy of this invention after oil cooling or furnace cooling + air cooling is higher than that of the FGH4097 alloy.

[0183] In summary, the mechanical property test results show that the overall properties of the alloy of this invention, such as tensile strength, creep resistance, impact resistance, and fatigue resistance, are better than those of the FGH4097 alloy. In particular, the tensile strength and high-temperature creep resistance of the alloy of this invention are higher than those of FGH4097.

[0184] Finally, it should be noted that the above embodiments are only used to illustrate the technical solutions of the present invention, and not to limit them; although the present invention has been described in detail with reference to the foregoing embodiments, those skilled in the art should understand that modifications can still be made to the technical solutions described in the foregoing embodiments, or equivalent substitutions can be made to some or all of the technical features; and these modifications or substitutions do not cause the essence of the corresponding technical solutions to deviate from the scope of the technical solutions of the embodiments of the present invention.

Claims

1. A nickel-based superalloy, characterized in that, By weight percentage, it includes the following components: C 0.05%~0.063%, Co 15.0%~15.04%, Cr 10.08%~10.13%, W 5.74%~5.81%, Mo 3.30%~3.34%, Al 3.73%~3.78%, Ti 3.68%~3.73%, Nb 1.77%~1.81%, Hf 0.20%~0.40%, B 0.010%~0.014%, Zr 0.008%~0.01%, Mg 0.002%~0.006%, balance Ni; The nickel-based superalloy comprises a γ phase, a γ' phase, carbides, borides, and a TCP phase; The average grain size of the nickel-based superalloy is 25~35μm; The mass percentage of the γ′ phase in the nickel-based superalloy is 57%~59%; The nickel-based superalloy has an average size of 0.5~4 μm for the primary γ′ phase, an average size of 0.2~0.8 μm for the secondary γ′ phase, and an average size of 0.04~0.07 μm for the tertiary γ′ phase. The nickel-based superalloy includes at least one of the following features (1) to (3); (1) The room temperature tensile strength of the nickel-based superalloy is 1510~1590MPa, the room temperature yield strength is 1020~1185MPa, the room temperature elongation after fracture is 14%~25%, and the room temperature reduction of area is 17%~23%; (2) The nickel-based high-temperature alloy has a tensile strength of 1090~1250MPa, a yield strength of 913~1050MPa, an elongation after fracture of 12%~35%, and a reduction of area of ​​12%~32% at 750℃; (3) The nickel-based high-temperature alloy has a composite creep rupture time of 180~220h at 650℃ and 990MPa, and a composite creep rupture time of 190~210h at 750℃ and 637MPa.

2. The nickel-based superalloy according to claim 1, characterized in that, In the nickel-based superalloy, the mass ratio of Ti to Al is (0.97~1):

1.

3. The nickel-based superalloy according to claim 1, characterized in that, In the nickel-based superalloy, the total mass percentage of Al, Ti, Nb, and Hf is 9.38% to 9.71%.

4. The method for preparing the nickel-based superalloy according to any one of claims 1 to 3, characterized in that, Includes the following steps: Alloy powder is prepared by mixing ingredients according to the specified ratio; the alloy powder is then subjected to hot isostatic pressing, solution treatment and aging treatment in sequence to obtain the nickel-based high-temperature alloy.

5. The method for preparing the nickel-based superalloy according to claim 4, characterized in that, The hot isostatic pressing process includes: heat treatment at 1190~1220℃ and 130~150MPa for 3~5 hours.

6. The method for preparing the nickel-based superalloy according to claim 4, characterized in that, The solution treatment includes: holding at 1170~1220℃ for 2~5 hours, followed by cooling.

7. The method for preparing the nickel-based superalloy according to claim 6, characterized in that, The cooling method for the solution treatment includes furnace cooling followed by air cooling; or, the cooling method for the solution treatment includes oil cooling. And / or, the cooling rate of the solution treatment is 30~40℃ / min; or, the cooling rate of the solution treatment is ≥80℃ / min.

8. The method for preparing the nickel-based superalloy according to claim 4, characterized in that, The aging treatment includes: heat treatment at 870±10℃ for 6~9h, followed by air cooling, heat treatment at 760±10℃ for 14~18h, followed by air cooling.

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