Nickel-based high-temperature alloy and preparation method thereof

Abstract

The invention relates to the technical field of alloys, in particular to a nickel-based high-temperature alloy and a preparation method thereof. The invention provides a nickel-based high-temperature alloy which comprises the following components in percentage by mass: 0.04% to 0.06% of C, 15.0% to 15.04% of Co, 10.08% to 10.13% of Cr, 5.74% to 5.81% of W, 3.30% to 3.34% of Mo, 3.73% to 3.78% of Al, 3.68% to 3.73% of Ti, 1.77% to 1.81% of Nb, 0.20% to 0.40% of Hf, 0.010% to 0.014% of B, 0.008% to 0.01% of Zr, 0.002% to 0.006% of Mg and the balance of Ni. Compared with an FGH4097 alloy, the nickel-based high-temperature alloy disclosed by the invention has more excellent comprehensive performance by adjusting components.

Description

Technical Field

[0001] The present invention relates to the technical field of alloys, and in particular to a nickel-based high-temperature alloy and a preparation method thereof. Background Art

[0002] FGH4097 alloy is a precipitation-strengthened nickel-based powder high-temperature alloy with a γ′ phase mass fraction of about 61%. It exhibits excellent performance in high-temperature environments, especially in the temperature range of 650°C to 750°C. Its comprehensive mechanical properties are particularly outstanding, making it an ideal material for the preparation of key hot-end components in high-end equipment such as aircraft engines.

[0003] Direct hot isostatic pressing near-net forming technology is an efficient and economical process for preparing powder high-temperature alloy discs. The FGH4097 alloy discs, shafts, ring parts and other parts prepared by this technology not only have uniform structure, but also good performance consistency, which greatly improves the reliability and service life of the product. The advantage of this process is that it can reduce material waste, improve production efficiency, and ensure product quality. It is an ideal choice for preparing high-end powder high-temperature alloy parts.

[0004] At present, FGH4097 alloy has been widely used in key components such as turbine disks and compressor disks of aircraft engines and ground-based gas turbines. These components are subjected to extremely high temperatures and pressures in the engine, so the material is required to have excellent high temperature resistance, high pressure resistance and fatigue resistance. With the continuous development of high thrust-to-weight ratio aircraft engine technology, the operating temperature and performance requirements for key hot end components of the engine, such as turbine disks, are also constantly increasing. In order to meet the needs of higher performance engines, it is necessary to conduct more in-depth research and improvement on FGH4097 alloy to further improve its high temperature resistance, high pressure resistance and fatigue resistance, and provide stronger support for the development of high-end equipment such as aircraft engines.

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

[0006] The first object of the present invention is to provide a nickel-based high-temperature alloy which has higher tensile strength and high-temperature durability than FGH4097 alloy by adjusting the alloy composition and has excellent comprehensive performance.

[0007] The second object of the present invention is to provide a method for preparing the above-mentioned nickel-based high-temperature alloy, thereby improving the tensile strength and high-temperature durability of the nickel-based high-temperature alloy.

[0008] In order to achieve the above-mentioned purpose of the present invention, the following technical solutions are particularly adopted:

[0009] In a first aspect, the present invention provides a nickel-based high-temperature alloy, comprising the following components in terms of mass percentage:

[0010] C 0.05%~0.063%, Co 15.0%~15.04%, Cr 10.08%~10.13%, W5.74%~5.81%, Mo 3.30%~3.34%, Al 3.73%~3.78%, Ti 3.68%~3.73%, Nb1.77%~1.81%, Hf 0.20%~0.40%, B 0.010%~0.014%, Zr 0.008%~0.01%, Mg 0.002%~0.006%, and the balance is Ni.

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

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

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

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

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

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

[0017] (4) The average size of the primary γ′ phase of the nickel-based high-temperature alloy 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] Further, it includes at least one of the following features (1) to (3);

[0019] (1) The room temperature tensile strength of the nickel-based high temperature alloy 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 cross-sectional shrinkage is 17%-23%;

[0020] (2) The nickel-based high-temperature alloy has a tensile strength of 1090 to 1250 MPa, a yield strength of 913 to 1050 MPa, an elongation after fracture of 12% to 35%, and a cross-sectional shrinkage of 12% to 32% at 750°C;

[0021] (3) The composite endurance fracture time of the nickel-based high-temperature alloy at 650°C and 990MPa is 180 to 220 hours, and the composite endurance fracture time at 750°C and 637MPa is 190 to 210 hours.

[0022] In a second aspect, the present invention further provides a method for preparing the nickel-based high-temperature alloy as described above, comprising the following steps:

[0023] The alloy powder is prepared after the ingredients are mixed according to the ratio; the alloy powder is subjected to hot isostatic pressing treatment, solution treatment and aging treatment in sequence to obtain the nickel-based high-temperature alloy.

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

[0025] Furthermore, the solution treatment comprises: heat preservation treatment at 1170-1220° C. for 2-5 hours and then cooling.

[0026] Further, the cooling method of the solid solution treatment includes furnace cooling followed by air cooling; or, the cooling method of the solid 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 preservation treatment at 870±10°C for 6 to 9 hours, air cooling, heat preservation treatment at 760±10°C for 14 to 18 hours, and air cooling.

[0029] Compared with the prior art, the present invention has the following beneficial effects:

[0030] 1. The nickel-based high-temperature alloy provided by the present invention adjusts the alloy composition on the basis of the FGH4097 alloy composition, improves the solid solution strengthening effect by increasing the C, Cr, and W contents and reducing the Co and Mo contents; increases the Ti content and reduces the Al and Nb contents, and increases the complete solid solution temperature of the γ′ phase under the condition that the total amount of the γ′ phase-forming elements Al, Ti, Nb, and Hf remains basically unchanged; thereby, the comprehensive properties of the nickel-based high-temperature alloy, such as tensile properties, endurance properties, impact properties, and fatigue properties, are better than those of the FGH4097 alloy, especially the tensile strength and high-temperature endurance properties are better than those of the FGH4097 alloy.

[0031] 2. The preparation method of the nickel-based high-temperature alloy of the present invention improves the performance of the nickel-based high-temperature alloy by studying the hot isostatic pressing and heat treatment processes. The room temperature tensile strength, high temperature tensile strength and endurance strength of the prepared nickel-based high-temperature alloy are better than those of the FGH4097 alloy prepared by direct hot isostatic pressing and near-net forming. BRIEF DESCRIPTION OF THE DRAWINGS

[0032] In order to more clearly illustrate the technical solutions of the embodiments of the present invention, the drawings required for use in the embodiments will be briefly introduced below. It should be understood that the following drawings only show certain embodiments of the present invention and therefore should not be regarded as limiting the scope. For ordinary technicians in this field, other related drawings can be obtained based on these drawings without paying creative work.

[0033] Figure 1 This is the grain structure of the small ingot according to Example 1 of the present invention.

[0034] Figure 2 This is the grain structure of the small ingot according to Example 2 of the present invention.

[0035] Figure 3 This is the grain structure of the small ingot according to Example 3 of the present invention.

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

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

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

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

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

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

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

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

[0044] Fig.12 This is the γ′ phase characteristic of the DP-HT-OC-1 small ingot of the present invention.

[0045] Fig.13 This is the γ′ phase characteristic of the DP-HT-OC-3 small ingot of the present invention.

[0046] Fig.14This is the γ′ phase characteristic of the DP-HT-OC-9 small ingot of the present invention. DETAILED DESCRIPTION

[0047] The technical scheme of the present invention will be clearly and completely described below in conjunction with the accompanying drawings and specific embodiments, but it will be understood by those skilled in the art that the following described embodiments are part of the present invention and are only used to illustrate the present invention and should not be considered as limiting the scope of the present invention. If specific conditions are not specified in the embodiments, they are carried out according to normal conditions or conditions recommended by the manufacturer. If the manufacturer is not specified for the reagents or instruments used, they are all conventional products that can be purchased commercially.

[0048] In some embodiments of the present invention, a nickel-based high-temperature alloy is provided, which comprises the following components in terms of mass percentage:

[0049] C 0.05%~0.063%, Co15.0%~15.04%, Cr10.08%~10.13%, W5.74%~5.81%, Mo3.30%~3.34%, Al3.73%~3.78%, Ti3.68%~3.73%, Nb1.77%~1.81%, Hf0.20%~0.40%, B0.010%~0.014%, Zr0.008%~0.01%, Mg0.002%~0.006%, and the balance is Ni.

[0050] In different embodiments, the mass percentages of the components in the nickel-based high-temperature alloy may 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 range between two thereof; the mass percentage of Co can be 15.0%, 15.01%, 15.02%, 15.03%, 15.04% or any range between two thereof; the mass percentage of Cr can be 10.08%, 10.09%, 10.10%, 10.11%, 10.12%, 10.13% or any range between two thereof; the mass percentage of W can be 5.74%, 5.7 The mass percentage of Mo may be 3.30%, 3.31%, 3.32%, 3.33%, 3.34% or a range between any two thereof; the mass percentage of Al may be 3.73%, 3.74%, 3.75%, 3.76%, 3.77%, 3.78% or a range between any two thereof; the mass percentage of Ti may be 3.68%, 3.69%, 3.70%, 3.71%, 3.72%, 3.73% or a range between any two thereof. The mass percentage of Nb can be 1.77%, 1.78%, 1.79%, 1.80%, 1.81% or a range between any two thereof; 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 a range between any two thereof, 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 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 a range between any two thereof, preferably 0.0086% to 0.01%; the mass percentage of Mg can be 0.002%, 0.003%, 0.004%, 0.005%, 0.006% or a range between any two thereof.

[0052] In some embodiments, in the nickel-based high-temperature alloy, the mass ratio of Ti to Al is (0.97-1):1; typically but not limiting, for example, the mass ratio of Ti to Al can be 0.97:1, 0.98:1, 0.99:1, 1:1 or a range of any two thereof.

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

[0054] The nickel-based high-temperature alloy of the present invention is based on the composition of FGH4097 alloy, and the solid solution strengthening effect is improved by increasing the contents of C, Cr and W and reducing the contents of Co and Mo; the mass ratio of Ti to Al is made close to 1:1 by increasing the content of Ti and reducing the contents of Al and Nb; the total mass percentage of the main constituent elements of the γ′ phase, Al, Ti, Nb and Hf, is 9.38%-9.71%, and 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, the complete dissolution temperature of the γ′ phase is 1188-1205°C, and the mass percentage of the γ′ phase is 57%-59%.

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

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

[0057] The equilibrium phases in nickel-based high-temperature alloys include matrix γ phase, γ' phase, carbides (MC, M 6 C.M. 23 C 6 ), boride (M 3 B 2 、MB 2 ) and trace amounts of TCP phase (μ phase, σ phase);

[0058] The average grain size of the nickel-based high-temperature alloy is 25 to 35 μm (for example, 25 μm, 27 μm, 29 μm, 31 μm, 33 μm, 55 μm, etc.); the mass percentage of the γ′ phase in the nickel-based high-temperature alloy is 57% to 59%;

[0059] The average size of the primary γ′ phase of the nickel-based high-temperature alloy is 0.5 to 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.), the average size of the secondary γ′ phase is 0.2 to 0.8 μm (e.g., 0.2 μm, 0.4 μm, 0.6 μm, 0.8 μm, etc.), and the average size of the tertiary γ′ phase is 0.04 to 0.07 μm (e.g., 0.04 μm, 0.05 μm, 0.06 μm, 0.07 μm, etc.). In some embodiments, the properties of the nickel-based high-temperature alloy are as follows:

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

[0061] High temperature (750°C) tensile properties of nickel-based high-temperature alloys: tensile strength of 1090-1250MPa (e.g., 1090MPa, 1110MPa, 1130MPa, 1150MPa, 1170MPa, 1190MPa, 1210MPa, 1230MPa, 1250MPa, etc.), yield strength of 913-1050MPa (e.g., 913MPa, 1000MPa, 1020MPa, 1040MPa, 1050MPa, etc.), elongation after fracture of 12%-35% (e.g., 12%, 15%, 20%, 25%, 30%, 35%, etc.), and cross-sectional shrinkage of 12%-32% (e.g., 12%, 15%, 20%, 25%, 30%, 32%, etc.);

[0062] Room temperature impact properties of nickel-based high-temperature alloys: absorbed energy is 27-42J (for example, 27J, 29J, 31J, 33J, 35J, 37J, 39J, 41J, 42J, etc.), and impact toughness is 33-53J / cm 2 (For example, 33 J / cm 2 、35J / cm 2 、37J / cm 2、39J / 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 performance of nickel-based high temperature alloys: fatigue life at 650°C and 1100MPa is 11000~50000N (11000N, 20000N, 30000N, 40000N, 50000N, etc.);

[0064] Smooth and durable performance of nickel-based high-temperature alloy at 650°C and 1140MPa: fracture time is 7 to 61h (for example, 7h, 10h, 20h, 30h, 40h, 50h, 60h, etc.), elongation after fracture is 8% to 27% (for example, 8%, 10%, 14%, 18%, 22%, 25%, 27%, etc.), and section reduction is 8% to 24% (for example, 8%, 10%, 14%, 18%, 22%, 24%, etc.);

[0065] Smooth and durable performance of nickel-based high-temperature alloy at 750°C and 750MPa: fracture time is 29 to 130h (for example, 29h, 50h, 70h, 90h, 110h, 130h, etc.), elongation after fracture is 4.5% to 15% (for example, 4.5%, 6%, 8%, 10%, 12%, 14%, 15%, etc.), and section shrinkage is 5% to 18% (for example, 5%, 7%, 9%, 11%, 13%, 15%, 17%, 18%, etc.);

[0066] Composite endurance properties of nickel-based high-temperature alloys: the composite endurance fracture time at 650°C and 990MPa is 180 to 220h (for example, 180h, 200h, 220h, etc.), and the composite endurance fracture time at 750°C and 637MPa is 190 to 210h (for example, 190h, 200h, 210h, etc.).

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

[0068] The alloy powder is prepared after the ingredients are mixed according to the proportion; the alloy powder is subjected to hot isostatic pressing treatment, solid 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: after mixing the ingredients according to the ratio, vacuum horizontal continuous casting (VHCC) is used to prepare a master alloy rod; the master alloy rod is atomized and pulverized by a plasma rotating electrode method (PREP) to obtain the alloy powder.

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

[0071] In some embodiments, the alloy powder is sequentially screened, electrostatically separated, and powder inspected, and then placed in a package for degassing, casing, and sealing, and then subjected to hot isostatic pressing.

[0072] In some embodiments, the hot isostatic pressing treatment includes: heat preservation treatment at 1190-1220°C and 130-150MPa for 3-5h; typically but not limiting, for example, the temperature of the hot isostatic pressing treatment can be 1190°C, 1200°C, 1210°C, 1220°C or a range of any two thereof; preferably, the temperature of the hot isostatic pressing treatment is 1200°C and the time of the hot isostatic pressing treatment is 4h. The present invention adopts the above hot isostatic pressing treatment to obtain a nickel-based high-temperature alloy with better γ′ phase distribution, quantity, size and morphology.

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

[0074] In some embodiments, the cooling method of the solution treatment includes furnace cooling and then air cooling; or, the cooling method of the solution treatment includes oil cooling.

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

[0076] The cooling method and cooling rate of the solution treatment have a significant effect on the grain size and performance of the obtained nickel-based high-temperature alloy. The use of the above cooling method and cooling rate can further improve the high-temperature durability of the alloy, increase the durability life and solve the problem of composite durability notch sensitivity.

[0077] In some embodiments, the aging treatment comprises: heat preservation treatment at 870±10°C for 6-9 hours, air cooling, heat preservation treatment at 760±10°C for 14-18 hours, and air cooling. Preferably, the aging treatment of the present invention is a two-stage aging treatment of heat preservation treatment at 870°C for 8 hours, heat preservation treatment at 760°C for 16 hours, and air cooling.

[0078] The microstructure and properties of the nickel-based high-temperature alloy obtained by the alloy powder with a particle size of 50-100 μm after hot isostatic pressing at 1200°C, solution treatment (temperature of 1180-1210°C, cooling method of 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] The room temperature (25°C) tensile properties of nickel-based high-temperature alloys are as follows: tensile strength is 1514-1557MPa, yield strength is 1024-1068MPa, elongation after fracture is 19.5%-24.5%, and section shrinkage is 20%-23%;

[0081] High temperature (750℃) tensile properties of nickel-based high temperature alloys: tensile strength is 1090~1160MPa, yield strength is 913~1000MPa, elongation after fracture is 23%~35%, and section shrinkage is 17%~32%;

[0082] Composite endurance properties of nickel-based high-temperature alloys: The composite endurance fracture time at 650°C and 990MPa is 210-220h, and the composite endurance fracture time at 750°C and 637MPa is 190-200h.

[0083] The microstructure and properties of the nickel-based high-temperature alloy obtained by the alloy powder with a particle size of 50-100 μm after hot isostatic pressing at 1200°C, solution treatment (temperature of 1180-1210°C, cooling method of 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] The room temperature (25°C) tensile properties of nickel-based high-temperature alloys are as follows: tensile strength is 1556-1587MPa, yield strength is 1129-1185MPa, elongation after fracture is 14.5%-17.5%, and section shrinkage is 17-18%;

[0086] High temperature (750℃) tensile properties of nickel-based high temperature alloys: tensile strength is 1200-1250MPa, yield strength is 990-1050MPa, elongation after fracture is 12%-24%, and section shrinkage is 13%-25.5%;

[0087] Composite endurance properties of nickel-based high-temperature alloys: the composite endurance fracture time at 650°C and 990MPa is 180-190h, and the composite endurance fracture time at 750°C and 637MPa is 200-210h.

[0088] The microstructure and properties of the nickel-based high-temperature alloy obtained by the alloy powder with a particle size of ≤75μm after hot isostatic pressing at 1200℃, solution treatment (temperature of 1180-1210℃, cooling method of 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] The room temperature (25°C) tensile properties of nickel-based high-temperature alloys are: tensile strength of 1563-1568MPa, yield strength of 1138-1141MPa, elongation after fracture of 18.5%-19.5%, and section shrinkage of 19%-21%.

[0092] High temperature (750℃) tensile properties of nickel-based high temperature alloys: tensile strength of 1220-1250MPa, yield strength of 1030-1050MPa, elongation after fracture of 12%-19%, and reduction of area of ​​12%-17%. The preparation method of the nickel-based high temperature alloy of the present invention improves the performance of the nickel-based high temperature alloy by studying the hot isostatic pressing and heat treatment processes; the room temperature tensile strength, high temperature tensile strength and endurance strength of the prepared nickel-based high temperature alloy are better than those of the FGH4097 alloy prepared by direct hot isostatic pressing near net forming; it can meet the requirements of aircraft engines for 750℃ disk materials, and can be applied to different types of aircraft engines and high-end parts of ground combustion engines such as disks, shafts, and rings.

[0093] The low-cost, high-performance, high-reliability, high-corrosion-resistant and long-life nickel-based high-temperature alloy prepared by the present invention can be specifically 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, warning / reconnaissance helicopters, anti-ship helicopters, anti-mine helicopters, search and rescue helicopters, etc., as well as gas turbines used for ground power generation or other industrial fields, including aircraft engines and aircraft-derivative engines.

[0094] Embodiments 1 to 4

[0095] The preparation method of the nickel-based high-temperature alloy provided in this embodiment, the composition (wt %) of the prepared nickel-based high-temperature alloy is 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 margin 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 margin 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 margin 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 margin

[0098] Specifically, the preparation method of the nickel-based high-temperature alloy comprises the following steps:

[0099] After the ingredients are prepared according to the proportion, vacuum horizontal continuous casting (VHCC) is used to prepare the master alloy bar; the master alloy bar is subjected to atomization powder preparation (PREP) by plasma rotating electrode method to obtain alloy powder with a particle size of 50 to 100 μm; the alloy powder is subjected to hot isostatic pressing treatment and then the coating is removed, the alloy powder is divided into four parts in a longitudinal cross shape, and cut into 1 / 4 pieces for solution treatment and aging treatment in sequence to obtain small ingots;

[0100] The hot isostatic pressing treatment is as follows: heat preservation treatment at 1200°C and 140MPa for 4h;

[0101] The solution treatment is: heat preservation treatment at 1180-1210°C for 4 hours, cooling; the cooling method is furnace cooling and then air cooling; the cooling rate is 30-40°C / min; the temperatures of the solution treatments in Examples 1 to 4 are 1180°C, 1190°C, 1200°C and 1210°C, respectively; the aging treatment includes: heat preservation treatment at 870°C for 8 hours, air cooling, heat preservation treatment at 760°C for 16 hours, and air cooling.

[0102] Comparative Examples 1-2

[0103] The preparation process of the nickel-based high-temperature alloy provided in this comparative example, the composition (wt %) of the prepared nickel-based high-temperature alloy is 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 - margin 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 margin

[0106] The preparation method of the above nickel-based high-temperature alloy is referred to Example 1.

[0107] Test Example 1

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

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

[0110] Table 3

[0111]

[0112] from Figure 1 to Figure 4 As shown in Table 3, the nickel-based high-temperature alloys in Examples 1 to 4 gradually increase in grain size with increasing solution temperature (1180 to 1210°C) after the same aging heat treatment, with an average grain size of 28.32 to 34.95 μm and a grain size of 6.0 to 7.0. When the solution temperature is 1180°C, recrystallization is insufficient, the grain size distribution is uneven, and it is bimodal, with a small amount of PPB and some residual dendrites. The grain sizes of the nickel-based high-temperature alloys in Comparative Examples 1 to 2 are all 6.0 to 6.5.

[0113] from Figures 5 to 8As shown in Table 3, under the same solution holding time, cooling method and two-stage aging, the nickel-based high-temperature alloys in Examples 1 to 4 have little change in size of the primary γ′ phase after aging treatment as the solution temperature increases (1180-1210°C), with an average size of 2.58-2.63 μm, mainly in the form of irregular long strips, distributed on the grain boundaries. When the solution temperature is 1180℃, the number of primary γ′ phases on the grain boundary is relatively large; as the solution temperature increases, the number of primary γ′ phases gradually decreases, and the size of the secondary γ′ phase gradually increases, with an average size of 0.61-0.70μm; when the solution temperature is 1180℃, the secondary γ′ phase is mainly triangular and butterfly-shaped, and when the solution temperature is 1190-1210℃, the secondary γ′ phase is mainly cubic and field-shaped, and is undergoing splitting and distributed in the crystal; the size of the tertiary γ′ phase is not much different, with an average size of 0.07μm, all in small spherical shapes, mainly distributed between the primary γ′ phase and the secondary γ′ phase. In summary, the small ingot billet treated at 1200℃ for 4h has a better effect. The size of the γ′ phase in the nickel-based high-temperature alloy in Comparative Example 1 is finer than that in Examples 1-4, and the size of the γ′ phase in the nickel-based high-temperature alloy in Comparative Example 2 is close to or slightly coarser than that in Examples 1-4.

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

[0115] Table 4

[0116]

[0117] From the room temperature tensile properties in Table 4, it can be seen that the nickel-based high-temperature alloys in Examples 1 to 4, when the solution temperature is 1180°C (Example 1), the room temperature tensile strength of the alloy is relatively high, with an average value of 1547MPa, but the yield strength is slightly lower, with an average value of 1028MPa. Comprehensively comparing the strength and plasticity of the alloy, when the solution temperature is 1200-1210°C (Examples 2-4), the room temperature tensile strength of the alloy is relatively high, with an average value of 1532MPa and 1532MPa, and the yield strength is relatively high, with an average value of 1062MPa and 1030MPa, respectively. The average values ​​of elongation after fracture are 24% and 24%, respectively, and the average values ​​of cross-sectional shrinkage are 22% and 23%, respectively. Compared with Comparative Example 1, the room temperature tensile strength and plasticity of the nickel-based high-temperature alloys in Examples 1 to 4 are both higher; compared with Comparative Example 2, the strength is slightly lower, but the plasticity is higher.

[0118] From the high temperature tensile properties in Table 4, it can be seen that the strength and plasticity of the nickel-based high temperature alloys in Examples 1 to 4 increase with the increase of the solid solution temperature (1180-1210°C). When the solid solution temperature is 1200-1210°C, the tensile strength is 1113MPa and 1145MPa, the yield strength is 924MPa and 1000MPa, the elongation is 34% and 32%, and the cross-sectional shrinkage is 29% and 32%. Compared with the comparative example 1, the nickel-based high temperature alloys of Examples 3 to 4 have similar tensile strength and plasticity at 750°C.

[0119] Table 5

[0120]

[0121] As shown in Table 5, when the solution temperature is 1180°C, the impact energy and impact toughness of the nickel-based high-temperature alloys in Examples 1 to 4 are the lowest, with average values ​​of 35 J and 43 J / cm2, respectively. 2 When the solution temperature is 1190-1210℃, the impact energy and impact toughness of the alloy are relatively high, with the average values ​​ranging from 38 to 42 J and 48 to 52 J / cm 2 Compared with comparative examples 1 to 2, the room temperature impact energy and impact toughness of the nickel-based high temperature alloys in Examples 1 to 4 are higher than those in Comparative examples 1 to 2.

[0122] From the low cycle fatigue performance in Table 5, it can be seen that the low cycle fatigue life of the nickel-based high temperature alloys in Examples 1 to 4 at 650°C / 1100MPa is 11119 to 33012 cycles, wherein when the solution temperature is 1200°C, the longest low cycle fatigue life is 33012 cycles. Comparative Examples 1 to 2 are low cycle fatigue lives of 63000 cycles and 16400 cycles under different conditions, namely 650°C / 980MPa and 540°C / strain ratio 0.95, respectively.

[0123] Table 6

[0124]

[0125]

[0126] From the smooth and durable performance at 650℃ / 1140MPa in Table 6, it can be seen that the nickel-based high-temperature alloys in Examples 1 to 4 have little change in durability and plasticity with the increase of the solid solution temperature (1180-1210℃). When the solid solution temperature is 1200℃, the durability is as high as 9.74h, and the elongation and cross-sectional shrinkage are 25% and 23% respectively, which reduces the stress. The smooth and durable lives of Comparative Examples 1 to 2 at 650℃ / 1120MPa and 650℃ / 1035MPa are 10h and 80h respectively, and the plasticity of both is relatively low.

[0127] From the smooth and durable performance at 750°C / 750MPa in Table 6, it can be seen that the durable life of the nickel-based high-temperature alloys in Examples 1 to 4 is extended with the increase of the solid solution temperature (1180-1210°C). When the solid solution temperature is 1200-1210°C, the durable 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] From the composite stress resistance performance at 650°C / 1140MPa in Table 7, it can be seen that the sustained fracture life of the nickel-based high-temperature alloy in Example 3 is 98h higher than that in Comparative Example 1, and the plasticity is equivalent; from the composite stress resistance performance at 750°C / 750MPa in Table 7, it can be seen that the sustained fracture life of the nickel-based high-temperature alloy in Example 3 is 65h higher than that in Comparative Example 1, and the plasticity is equivalent.

[0131] Test Example 2

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

[0133] The master alloy rod was prepared by vacuum horizontal continuous casting technology (VHCC), and the nickel-based high-temperature alloy powder of different particle sizes was prepared by plasma rotating electrode method (PREP). The composition of the nickel-based high-temperature alloy powder was the same as that in Example 3. After being put into a can for degassing, canning and sealing, the grain size and grain grade statistics after hot isostatic pressing at different temperatures of 1210°C and 1200°C and 130 MPa for 4 hours are shown in Table 8, and the γ′ phase characteristics are shown in Table 9.

[0134] Table 8

[0135]

[0136] As shown in Table 8, the average grain size of the alloy powder with a particle size of 50-100 μm after hot isostatic pressing at 1200℃ and 1210℃ was 31.74-32.53 μm, the grain size was 6.5-7.0, the grain size distribution was relatively uniform, the recrystallization was relatively sufficient, and there were a small amount of primary particle boundaries (PPB) and residual dendrites. The average grain size of the alloy powder with a particle size of ≤75 μm after hot isostatic pressing at 1200℃ and 1210℃ was 28.42-29.49 μm, the grain size was 6.5-7.0, which was finer than that of the alloy powder with a particle size of 50-100 μm after hot isostatic pressing, and the distribution was relatively uniform, the recrystallization was relatively sufficient, and there were a small amount of PPB and residual dendrites.

[0137] Table 9

[0138]

[0139] As can be seen from Table 9, after the alloy powder with a particle size of 50-100μm was hot isostatically pressed at 1200℃ and 1210℃, the size of the primary γ′ phase did not change much, and the primary γ′ phase characteristics were all irregular long strips, but the amount of the primary γ′ phase after hot isostatic pressing at 1200℃ was more than that at 1210℃, and it was mainly distributed on the grain boundaries; the size of the secondary γ′ phase was slightly smaller than that at 1210℃, and the secondary γ′ phase characteristics were that the precipitation amount was large, mainly in the shape of cubes and butterflies, and splitting was occurring; the size of the tertiary γ′ phase was slightly larger than that at 1210℃, and the tertiary γ′ phase characteristics were all small spheres, and there were two types of tertiary γ′ phases, large and small. After hot isostatic pressing at 1200℃ and 1210℃, the size of the primary γ′ phase, secondary γ′ phase and tertiary γ′ phase of alloy powder with particle size ≤75μm did not change much. The primary γ′ phase was irregular and long, but the amount of primary γ′ phase after hot isostatic pressing at 1200℃ was more than that at 1210℃, and it was mainly distributed on the grain boundary. The amount of secondary γ′ phase precipitated was large, mainly in the shape of cube and butterfly, and was splitting. The tertiary γ′ phase was small spherical, and the size was divided into large and small tertiary γ′ phases. The size of the primary γ′ phase of alloy powder with particle size ≤75μm after hot isostatic pressing was slightly smaller than that of alloy powder with particle size of 50-100μm, the size of the secondary γ′ phase was slightly larger, and the size of the tertiary γ′ phase was equivalent. The distribution, quantity, size and morphology of the three sizes of γ′ phase of alloy powders with a particle size of 50-100 μm and a particle size of ≤75 μm after hot isostatic pressing at 1200°C are all better.

[0140] The BT-3 small ingot after removing the sheath was selected and divided into four in a longitudinal cross shape, and cut into 1 / 4 pieces for solution treatment and secondary aging treatment (after keeping at 870℃ for 8h, air cooling, keeping at 760℃ for 16h, air cooling), and then named as DP-HT-3 small ingot, with diameter and height dimensions of Φ80~90mm×h130~140mm.

[0141] The small ingots of BT-1, BT-3 and BT-9 after removing the sheath were selected and divided into four in a longitudinal cross shape, cut into 1 / 4 pieces for solution treatment (oil cooling) and secondary aging treatment (after keeping treatment at 870℃ for 8h, air cooling, keeping treatment at 760℃ for 16h, air cooling), and then named DP-HT-OC-1 small ingot, DP-HT-OC-3 small ingot and DP-HT-OC-9 small ingot, with diameter and height dimensions of Φ80~90mm×h130~140mm.

[0142] The grain size and grain grade 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 in Figure 2, the γ′ phase characteristics are as follows: Figure 7 As shown, it is the same as Example 3; the grain structure of the DP-HT-OC-1 small ingot, DP-HT-OC-3 small ingot and DP-HT-OC-9 small ingot is as shown Fig. 9 , Fig.10 , Fig.11 As shown ( Figures 9 to 11 The scale bar of Figure a is 100 μm, and the scale bar of Figure b is 20 μm); the γ′ phase characteristics of DP-HT-OC-1 small ingot, DP-HT-OC-3 small ingot and DP-HT-OC-9 small ingot at different magnifications are shown in Figure 1. Fig.12 , Fig.13 , Fig.14 shown.

[0143] Table 12

[0144]

[0145] from Figures 9 to 11 As shown in Table 12, the grain size of the DP-HT-OC-1 small ingot, DP-HT-OC-3 small ingot and DP-HT-OC-9 small ingot after solution treatment and oil cooling is smaller than that of the ingot after solution treatment and furnace cooling + air cooling. The grain size of the DP-HT-OC-1 small ingot after oil cooling is about 32.69μm, and the grain size is 6.5-7.0. The grain size after furnace cooling + air cooling is slightly coarser, about 34.25μm. The grain size of the DP-HT-OC-3 small ingot after oil cooling is about 32.92μm, and the grain size is 6.5-7.0. The grain size after furnace cooling + air cooling is about 33.67μm, and the grain size is 6.0-6.5. The grain size of DP-HT-OC-9 small ingot after oil cooling is about 25.62μm, and the grain size is 7.0-7.5. The grain size after furnace cooling + air cooling is about 29.15μm, and the grain size is 6.5-7.0. It can be seen that the solid solution cooling rate has a great influence on the grain size of small ingots, and oil cooling can reduce the grain size of the alloy.

[0146] Table 13

[0147]

[0148] from Figure 12 to Figure 14 As shown in Table 13, the size of the γ′ phase of DP-HT-OC-1, DP-HT-OC-3 and DP-HT-OC-9 small ingots using oil cooling is smaller than that of furnace cooling + air cooling. Taking the DP-HT-OC-3 small ingot using oil cooling and the DP-HT-3 using furnace cooling + air cooling as examples for comparison, it can be seen that the size of the primary γ′ phase and the secondary γ′ phase of the DP-HT-OC-3 small ingot using oil cooling are smaller than those of furnace cooling + air cooling, and the tertiary γ′ phase is close to or slightly smaller. After oil cooling, the size of the primary γ′ phase in the DP-HT-OC-3 small ingot is small and irregularly long, mainly distributed on the grain boundary. The secondary γ′ phase precipitates in large quantities and is small in size, mainly distributed in the crystal in a cubic shape. The tertiary γ′ phase precipitates in small quantities and is small in size, 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 are oil-cooled, and the γ′ phase characteristics are similar to those of the DP-HT-OC-3 small ingots using oil cooling.

[0149] FGH4097 standard A (alloy composition see "China High Temperature Alloy Standard Handbook (Volume 2)") heat treatment system: solution heat treatment is (1120~1160)±10℃ / (2~4)h / salt quenching or oil quenching, primary aging is 870℃±10℃ / 1h / AC, secondary aging is 650℃±10℃ / 24h / AC; FGH4097 standard B (alloy composition is the same as standard A) heat treatment system: solution heat treatment is (1140~1160)±10℃ / (1~2)h / salt quenching, primary aging is 870℃±10℃ / 1.5h / AC, secondary aging is 650℃±10℃ / 24h / AC.

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

[0151] Table 14

[0152]

[0153] As shown in Table 14, the room temperature tensile strength of the alloys treated with the same solution temperature of 1200°C, the same holding time of 4h, different solution cooling rates (furnace cooling + air cooling, oil cooling) + aging treatment is compared. The average room temperature tensile strength of the alloy treated with solution oil cooling is 40MPa higher than that of furnace cooling + air cooling, the average yield strength of the former is 127MPa higher than that of the latter, the plasticity of the former is slightly lower than that of the latter, the elongation after fracture of the former is 8% lower than that of the latter, and the section shrinkage is 6% lower than that of the latter. The room temperature tensile strength of the alloys treated with solution oil cooling of different powder sizes (50-100μm and ≤75μm) is compared. The strength and plasticity of the two are similar, with no significant difference. Overall, the room temperature tensile properties of the alloy of the present invention and the FGH4097 alloy (two aging systems) are compared. The room temperature tensile strength and plasticity of the alloy of the present invention are both higher than those of the FGH4097 alloy.

[0154] Table 15

[0155]

[0156]

[0157] From the high temperature tensile properties of the alloys in Table 15, it can be seen that the 750℃ tensile properties of the alloys after the same solution temperature of 1200℃, the same holding time of 4h, different solution cooling rates (furnace cooling + air cooling, oil cooling) + aging treatment, the average tensile strength of the alloy using solution oil cooling is 112MPa higher than that of furnace cooling + air cooling, the average yield strength of the former is 97MPa higher than that of the latter, the plasticity of the former is lower than that of the latter, the elongation after fracture of the former is 16% lower than that of the latter, and the section shrinkage rate of the former is 10% lower than that of the latter. Comparing the 750℃ high temperature tensile properties of the alloys with different powder particle sizes (50-100μm and ≤75μm) solution oil cooling, the strength and plasticity of the two are similar, with no obvious difference. Overall, the 750℃ tensile properties of the alloy of the present invention and the FGH4097 alloy (two aging systems) are compared. The high temperature tensile strength of the alloy of the present invention oil cooling is similar to that of the FGH4097 alloy, but the plasticity of the alloy of the present invention using oil cooling is lower than that of FGH4097. The high temperature tensile strength of the alloy using furnace cooling + air cooling of the present invention is lower than that of the FGH4097 alloy, and the plasticity of the alloy is slightly higher than or similar to that of the FGH4097 alloy.

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

[0159] Table 16

[0160]

[0161] As shown in Table 16, the impact energy and impact toughness of the alloys treated with the same solution temperature of 1200 °C, the same holding time of 4 h, different solution cooling rates (furnace cooling + air cooling, oil cooling) + aging treatment are compared. The average impact energy and impact toughness of the alloy treated with solution oil cooling are 7 J and 8 J / cm lower than those treated with furnace cooling + air cooling, respectively. 2 . Comparing the impact energy and impact toughness of alloys after solution oil cooling with different powder particle sizes (50-100μm and ≤75μm), the impact energy and impact toughness of the two are similar, with no significant difference. Overall comparison of the room temperature impact properties of the alloy of the present invention and the FGH4097 alloy (two aging systems), the room temperature impact energy and impact toughness of the alloy of the present invention using oil cooling are slightly lower than those of the FGH4097 alloy. The room temperature impact energy and impact toughness of the alloy of the present invention using furnace cooling + air cooling are similar to those of the 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] As can be seen from Table 17, the low-cycle fatigue performance of the alloys treated with the same solution temperature of 1200°C, the same holding time of 4h, different solution cooling rates (furnace cooling + air cooling, oil cooling) + aging treatment is compared. The low-cycle fatigue life of the alloy treated with solution oil cooling is 6791 cycles lower than that of furnace cooling + air cooling. The low-cycle fatigue performance of the alloys treated with solution oil cooling of different powder sizes (50-100μm and ≤75μm) is compared. The low-cycle fatigue life of the alloy with a powder size of ≤75μm is 14516 cycles higher than that of the alloy with a powder size of 50-100μm. Overall, the low-cycle fatigue performance of the alloy of the present invention and the FGH4097 alloy (two aging systems) is compared. The low-cycle fatigue life of the alloy of the present invention after oil cooling or furnace cooling + air cooling is slightly higher than that of the FGH4097 alloy.

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

[0167] Table 18

[0168]

[0169]

[0170] As shown in Table 18, the smooth and durable performance of the alloys treated with the same solution temperature of 1200℃, the same holding time of 4h, different solution cooling rates (furnace + air cooling and oil cooling) + aging treatment at 650℃ / 1140MPa is compared. The alloy with solution oil cooling has a 27.86h longer durability than that with furnace cooling + air cooling, and the former has lower plasticity than the latter. The elongation and reduction of section of the former are 18.5% and 14.8% lower than those of the latter, respectively. The smooth and durable performance of the alloys with different powder sizes (50-100μm and ≤75μm) after solution oil cooling is compared. The durability of the alloy with 50-100μm powder size is 19h longer than that of the alloy with ≤75μm powder size, but the plasticity of the former coarse powder is worse than that of the latter fine powder, and the elongation and reduction of section of the former are 5.5% and 3.8% lower than those of the latter, respectively. After reducing the smoothness and durability stress, the smoothness and durability performance at 650℃ / 1100MPa is compared with the smoothness and durability performance of the alloys treated with the same solution temperature of 1200℃, the same holding time of 4h, different solution cooling rates (furnace cooling + air cooling, oil cooling) + aging treatment at 650℃ / 1100MPa. After reducing the stress, the durability life of the alloys is extended. The durability life of the alloy using solution oil cooling is 61h, which is 42.9h higher than that of furnace cooling + air cooling. The plasticity of the former is lower than that of the latter, and the elongation after fracture and the section shrinkage rate of the former are 12.7% and 13.2% lower than those of the latter, respectively. By reducing the stress, the smoothness and durability performance of the alloy of the present invention and the FGH4097 alloy (primary aging system) at 650℃ / 1025-1120MPa are compared. The smoothness and durability life of the alloy of the present invention after oil cooling or furnace cooling + air cooling is higher than that of the FGH4097 alloy.

[0171] Table 19

[0172]

[0173]

[0174] As shown in Table 19, the smooth and durable performance of the alloys treated with the same solution temperature of 1200℃, the same holding time of 4h, different solution cooling rates (furnace + air cooling and oil cooling) + aging treatment at 750℃ / 750MPa is compared. The alloy with solution oil cooling has a 12.9h lower durability than that with furnace cooling + air cooling, and the former has lower plasticity than the latter. The elongation and area reduction of the former are 7.7% and 8.9% lower than those of the latter, respectively. The smooth and durable performance of the alloys with different powder sizes (50-100μm and ≤75μm) after solution oil cooling is compared. The durability and plasticity of the two are similar. After reducing the smoothness and durability stress, the smoothness and durability performance at 750℃ / 690MPa is compared with the smoothness and durability performance of the alloy at 750℃ / 690MPa after the same solution temperature of 1200℃, the same holding time of 4h, different solution cooling rates (furnace cooling + air cooling, oil cooling) + aging treatment. After reducing the stress, the durability of the alloy is extended. The durability of the alloy using oil cooling is 130h, which is similar to furnace cooling + air cooling, and the plasticity of the two is also similar. Reducing the stress to compare the smoothness and durability of the alloy of the present invention and the FGH4097 alloy (primary aging system) at 750℃ / 690MPa, the smoothness and durability of the alloy of the present invention after oil cooling or furnace cooling + air cooling is higher than that of the FGH4097 alloy.

[0175] The composite endurance performance of each small ingot was tested, and the results are shown in Tables 20 and 21.

[0176] Table 20

[0177]

[0178] As shown in Table 20, the durability and plasticity of the alloy using oil cooling are 32 hours lower than those after furnace cooling + air cooling, and its elongation after fracture and cross-sectional shrinkage are also 4.9% and 4.8% lower than those after furnace cooling + air cooling, respectively. Comparing the composite durability performance of the alloy of the present invention and the FGH4097 alloy (two aging systems) at 650°C / 980-990MPa, the composite durability of the alloy of the present invention after oil cooling or furnace cooling + air cooling is higher than that of the FGH4097 alloy.

[0179] Table 21

[0180]

[0181]

[0182] As can be seen from Table 21, the composite endurance performance of the alloy at 750°C / 637MPa after the same solution temperature of 1200°C, the same holding time of 4h, different solution cooling rates (furnace cooling + air cooling, oil cooling) + aging treatment is compared. The alloy endurance life of the oil-cooled alloy is 203.5h, which is similar to furnace cooling + air cooling. However, if the cooling rate of the oil cooling is too fast, the alloy is prone to fracture and notch phenomenon. Comparing the composite endurance performance of the alloy of the present invention and the FGH4097 alloy (primary aging system) at 750°C / 637MPa, the composite endurance life of the alloy of the present invention after oil cooling or furnace cooling + air cooling is higher than that of the FGH4097 alloy.

[0183] From the above mechanical properties test results, it can be seen that the comprehensive properties of the alloy of the present invention, such as tensile strength, endurance, impact and fatigue, are better than those of the FGH4097 alloy, especially the tensile strength and high temperature endurance of the alloy of the present 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, rather than to limit it. Although the present invention has been described in detail with reference to the aforementioned embodiments, those skilled in the art should understand that they can still modify the technical solutions described in the aforementioned embodiments, or replace some or all of the technical features therein with equivalents. However, these modifications or replacements 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 high-temperature alloy, characterized in that: According to mass percentage, it includes the following ingredients: 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%, Nb1.77%~1.81%, Hf0.20%~0.40%, B 0.010%~0.014%, Zr 0.008%~0.01%, Mg 0.002%~0.006%, and the balance is Ni.

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

1.

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

4. The nickel-based high-temperature alloy according to claim 1, characterized in that: The method comprises at least one of the following features (1) to (4); (1) The nickel-based high-temperature alloy includes a γ phase, a γ' phase, a carbide, a boride and a TCP phase; (2) The average grain size of the nickel-based high-temperature alloy is 25 to 35 μm; (3) The mass percentage of the γ′ phase in the nickel-based high-temperature alloy is 57% to 59%; (4) The average size of the primary γ′ phase of the nickel-based high-temperature alloy 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.

5. The nickel-based high-temperature alloy according to claim 1, characterized in that: The method comprises at least one of the following features (1) to (3); (1) The room temperature tensile strength of the nickel-based high temperature alloy 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 cross-sectional shrinkage is 17%-23%; (2) The nickel-based high-temperature alloy has a tensile strength of 1090 to 1250 MPa, a yield strength of 913 to 1050 MPa, an elongation after fracture of 12% to 35%, and a cross-sectional shrinkage of 12% to 32% at 750°C; (3) The composite endurance fracture time of the nickel-based high-temperature alloy at 650°C and 990MPa is 180 to 220 hours, and the composite endurance fracture time at 750°C and 637MPa is 190 to 210 hours.

6. The method for preparing the nickel-based high-temperature alloy according to any one of claims 1 to 5, characterized in that: The steps include: The alloy powder is prepared after the ingredients are mixed according to the ratio; the alloy powder is subjected to hot isostatic pressing treatment, solution treatment and aging treatment in sequence to obtain the nickel-based high-temperature alloy.

7. The method for preparing a nickel-based high-temperature alloy according to claim 6, characterized in that: The hot isostatic pressing treatment comprises: heat preservation treatment at 1190-1220° C. and 130-150 MPa for 3-5 hours.

8. The method for preparing a nickel-based high-temperature alloy according to claim 6, characterized in that: The solution treatment comprises: heat preservation treatment at 1170-1220° C. for 2-5 hours and then cooling.

9. The method for preparing a nickel-based high-temperature alloy according to claim 8, characterized in that: The cooling method of the solid solution treatment includes furnace cooling followed by air cooling; or the cooling method of the solid solution treatment includes oil cooling; 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.

10. The method for preparing a nickel-based high-temperature alloy according to claim 6, characterized in that: The aging treatment comprises: heat preservation treatment at 870±10° C. for 6 to 9 hours, air cooling, heat preservation treatment at 760±10° C. for 14 to 18 hours, and air cooling.

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