CN117926154A - Preparation method of nickel-based superalloy - Google Patents
Preparation method of nickel-based superalloy Download PDFInfo
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- CN117926154A CN117926154A CN202410121870.8A CN202410121870A CN117926154A CN 117926154 A CN117926154 A CN 117926154A CN 202410121870 A CN202410121870 A CN 202410121870A CN 117926154 A CN117926154 A CN 117926154A
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- PXHVJJICTQNCMI-UHFFFAOYSA-N Nickel Chemical compound [Ni] PXHVJJICTQNCMI-UHFFFAOYSA-N 0.000 title claims abstract description 86
- 229910000601 superalloy Inorganic materials 0.000 title claims abstract description 48
- 229910052759 nickel Inorganic materials 0.000 title claims abstract description 42
- 238000002360 preparation method Methods 0.000 title abstract description 21
- 238000001816 cooling Methods 0.000 claims abstract description 51
- 238000010438 heat treatment Methods 0.000 claims abstract description 42
- 238000005242 forging Methods 0.000 claims abstract description 36
- 238000000034 method Methods 0.000 claims abstract description 31
- 238000004321 preservation Methods 0.000 claims abstract description 30
- 238000007599 discharging Methods 0.000 claims abstract description 20
- 238000000265 homogenisation Methods 0.000 claims abstract description 18
- 238000003723 Smelting Methods 0.000 claims abstract description 8
- 230000000087 stabilizing effect Effects 0.000 claims abstract description 8
- 230000001550 time effect Effects 0.000 claims abstract description 4
- 230000008569 process Effects 0.000 claims description 13
- 229910052760 oxygen Inorganic materials 0.000 claims description 5
- 230000000630 rising effect Effects 0.000 claims description 5
- 229910052804 chromium Inorganic materials 0.000 claims description 4
- 229910052742 iron Inorganic materials 0.000 claims description 4
- 229910052750 molybdenum Inorganic materials 0.000 claims description 4
- 229910052698 phosphorus Inorganic materials 0.000 claims description 4
- 229910052726 zirconium Inorganic materials 0.000 claims description 4
- 229910001566 austenite Inorganic materials 0.000 claims description 3
- 238000004519 manufacturing process Methods 0.000 claims description 2
- 229910045601 alloy Inorganic materials 0.000 abstract description 82
- 239000000956 alloy Substances 0.000 abstract description 82
- 230000000052 comparative effect Effects 0.000 description 29
- 238000005728 strengthening Methods 0.000 description 13
- 239000006104 solid solution Substances 0.000 description 12
- 238000005336 cracking Methods 0.000 description 11
- 238000005204 segregation Methods 0.000 description 10
- 239000013078 crystal Substances 0.000 description 9
- 230000002045 lasting effect Effects 0.000 description 8
- 238000001556 precipitation Methods 0.000 description 8
- 230000009286 beneficial effect Effects 0.000 description 7
- 229910052799 carbon Inorganic materials 0.000 description 7
- 230000005496 eutectics Effects 0.000 description 7
- 239000011159 matrix material Substances 0.000 description 7
- 238000002844 melting Methods 0.000 description 7
- 239000000463 material Substances 0.000 description 6
- 230000008018 melting Effects 0.000 description 6
- 101000912561 Bos taurus Fibrinogen gamma-B chain Proteins 0.000 description 5
- 230000000694 effects Effects 0.000 description 5
- 239000000203 mixture Substances 0.000 description 5
- 230000009467 reduction Effects 0.000 description 5
- 239000000243 solution Substances 0.000 description 5
- 238000009826 distribution Methods 0.000 description 4
- 238000005096 rolling process Methods 0.000 description 4
- 230000035882 stress Effects 0.000 description 4
- 229910000831 Steel Inorganic materials 0.000 description 3
- 230000015572 biosynthetic process Effects 0.000 description 3
- 238000010586 diagram Methods 0.000 description 3
- 230000006872 improvement Effects 0.000 description 3
- 230000006698 induction Effects 0.000 description 3
- 150000001247 metal acetylides Chemical class 0.000 description 3
- 229910052757 nitrogen Inorganic materials 0.000 description 3
- 239000010959 steel Substances 0.000 description 3
- 239000000126 substance Substances 0.000 description 3
- 238000010313 vacuum arc remelting Methods 0.000 description 3
- UCKMPCXJQFINFW-UHFFFAOYSA-N Sulphide Chemical compound [S-2] UCKMPCXJQFINFW-UHFFFAOYSA-N 0.000 description 2
- 230000009471 action Effects 0.000 description 2
- 230000032683 aging Effects 0.000 description 2
- 238000005266 casting Methods 0.000 description 2
- 238000010622 cold drawing Methods 0.000 description 2
- 230000003111 delayed effect Effects 0.000 description 2
- 238000007711 solidification Methods 0.000 description 2
- 230000008023 solidification Effects 0.000 description 2
- 241001062472 Stokellia anisodon Species 0.000 description 1
- 230000004913 activation Effects 0.000 description 1
- 230000002411 adverse Effects 0.000 description 1
- 238000000137 annealing Methods 0.000 description 1
- 230000033228 biological regulation Effects 0.000 description 1
- 230000008859 change Effects 0.000 description 1
- 238000002485 combustion reaction Methods 0.000 description 1
- 230000007797 corrosion Effects 0.000 description 1
- 238000005260 corrosion Methods 0.000 description 1
- 210000001787 dendrite Anatomy 0.000 description 1
- 230000001627 detrimental effect Effects 0.000 description 1
- 238000009792 diffusion process Methods 0.000 description 1
- 238000004090 dissolution Methods 0.000 description 1
- 238000005324 grain boundary diffusion Methods 0.000 description 1
- 238000005098 hot rolling Methods 0.000 description 1
- 239000012535 impurity Substances 0.000 description 1
- 230000007774 longterm Effects 0.000 description 1
- 150000004767 nitrides Chemical class 0.000 description 1
- 230000006911 nucleation Effects 0.000 description 1
- 238000010899 nucleation Methods 0.000 description 1
- 230000003647 oxidation Effects 0.000 description 1
- 238000007254 oxidation reaction Methods 0.000 description 1
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- 229910052717 sulfur Inorganic materials 0.000 description 1
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Classifications
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
- C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
- C22F1/10—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of nickel or cobalt or alloys based thereon
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
- C21D9/28—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for plain shafts
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
- C21D9/40—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for rings; for bearing races
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C19/00—Alloys based on nickel or cobalt
- C22C19/03—Alloys based on nickel or cobalt based on nickel
- C22C19/05—Alloys based on nickel or cobalt based on nickel with chromium
- C22C19/051—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W
- C22C19/056—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W with the maximum Cr content being at least 10% but less than 20%
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
- C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
- C22F1/002—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working by rapid cooling or quenching; cooling agents used therefor
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- Chemical & Material Sciences (AREA)
- Engineering & Computer Science (AREA)
- Materials Engineering (AREA)
- Mechanical Engineering (AREA)
- Metallurgy (AREA)
- Organic Chemistry (AREA)
- Physics & Mathematics (AREA)
- Thermal Sciences (AREA)
- Crystallography & Structural Chemistry (AREA)
- Turbine Rotor Nozzle Sealing (AREA)
Abstract
The invention discloses a preparation method of a nickel-based superalloy, belongs to the technical field of superalloys, and solves the problem that the nickel-based superalloy in the prior art is difficult to simultaneously meet the comprehensive requirements of components on long service life of the alloy and low crack growth rate. Comprising the following steps: step 1: smelting to obtain an ingot; step 2: homogenizing heat treatment is carried out on the cast ingot; the homogenization heat treatment comprises 2 sections of heat preservation; 2, after the heat preservation of the 2 nd stage, furnace cooling to 1000+/-10 ℃ or lower, and then furnace discharging and air cooling; step 3: forging to prepare a bar blank and a forging; step 4: solution treatment; step 5: stabilizing and treating with time effect to obtain nickel-base superalloy. The nickel-based superalloy prepared by the method has high strength, good low cycle fatigue performance, low crack growth rate and excellent comprehensive performance.
Description
Technical Field
The invention relates to the technical field of high-temperature alloys, in particular to a preparation method of a nickel-based high-temperature alloy.
Background
At present, annular members and disk forgings for aviation and aerospace, such as turbine disks and blades of a combustion engine, need materials to ensure a certain lower crack expansion rate by considering high-temperature durability and creep fatigue at the same time, and have higher damage tolerance limit so as to ensure the service reliability of the materials; however, the conventional nickel-based alloy has a relatively high crack growth rate, and cannot meet the comprehensive requirements of some aerospace components on high alloy durability and low crack growth rate.
Disclosure of Invention
In view of the above, the present invention aims to provide a preparation method of a nickel-based superalloy, which is used for solving the problem that the existing nickel-based superalloy is difficult to simultaneously meet the comprehensive requirements of components on high alloy durability and low crack growth rate.
The aim of the invention is mainly realized by the following technical scheme:
The invention provides a preparation method of a nickel-based superalloy, which comprises the following steps:
step 1: smelting to obtain an ingot;
Step 2: homogenizing heat treatment is carried out on the cast ingot; the homogenization heat treatment comprises 2 sections of heat preservation; 2, after the heat preservation of the 2 nd stage, furnace cooling to 1000+/-10 ℃ or lower, and then furnace discharging and air cooling;
Step 3: forging to prepare a bar blank and a forging;
Step 4: solution treatment;
step 5: stabilizing and treating with time effect to obtain nickel-base superalloy.
Further, in step 2, the process steps of homogenizing heat treatment include:
S201, slowly heating the furnace temperature to 1165-1170 ℃ from the temperature of less than or equal to 400 ℃ for 10-15 h, and preserving heat for 35-37 h;
S202, continuously heating to 1200-1210 ℃, preserving heat for 38-42 h, cooling to below 1000+/-10 ℃, discharging and air cooling.
Further, in step 4, the solution treatment process is as follows: heating to 995-1020 ℃, preserving heat for 3-5 h, discharging and cooling.
In step 4, the temperature rising speed is 3-6 ℃/min.
In the step 4, the disc shaft parts with the wall thickness of 20-200 mm are cooled by discharging cooling oil; and (3) discharging air cooling is adopted for the annular piece with the wall thickness of 10-100 mm.
Further, in step 5, the stabilizing and time-efficient treatment process comprises: heating to 820-850 ℃ for heat preservation for 5-7 h, air cooling to 500+/-10 ℃, continuously heating to 740-770 ℃ for heat preservation for 8-10 h, and air cooling to room temperature.
In step 5, the temperature rising speed is controlled to be 2-3.5 ℃/min.
Further, the nickel-based superalloy comprises ：C：0.02％～0.06％,Cr：18.5％～20.0％,Co：13.0％～14.0％,Mo：4.0％～4.90％,Al：1.3％～1.6％,Ti：2.80％～3.25％,B：0.015％～0.03％,Zr：0.05％～0.15％,Ti/Al：2.25～2.38,(Al+Ti)：4.35％～4.58％,Fe：0.05％～2.0％,Mg≤0.005％,P：0.004％～0.009％,O：≤20PPm,N：≤20PPm,S≤10PPm, nickel in mass percent: the balance.
Further, the microstructure of the nickel-based superalloy obtained in the step 5 mainly comprises equiaxed austenite grains, boride and carbide which are uniformly dispersed and distributed, and gamma' phase which is dispersed and distributed in the grains.
Further, in the microstructure of the nickel-base superalloy, the carbide mainly includes M 23C6 and MC.
Compared with the prior art, the invention can at least realize one of the following beneficial effects:
a) In the preparation method of the nickel-based superalloy, the technological parameters of each step are precisely controlled, for example, homogenization heat treatment comprises 2 sections of heat preservation; the cracking tendency is reduced by furnace discharging air cooling after the 2 nd section is insulated and furnace cooling is carried out to below 1000+/-10 ℃; the combination of the accurate control of other steps and parameters of the steps avoids deformation, ensures that crystal grains of the alloy are uniform and tiny, ensures the best matching of the morphology and distribution of carbides in the alloy and the content and the size of other second phases, achieves the purposes of improving the strength of the alloy, reducing the cracking tendency and the crack expansion rate, and ensures the comprehensive performance of the nickel-based superalloy.
B) According to the preparation method of the nickel-based superalloy, the solid solution strengthening effect of the alloy and the grain boundary strength of the alloy are improved by accurately controlling the content of B, zr, fe, C, cr, co, al, ti and other single elements in the alloy; and by cooperatively controlling the values of B, zr and C, the best matching of the morphology and distribution of carbide in the alloy and the content and the size of other second phases is ensured, the risk crack sources in the alloy are reduced on the premise of not reducing the strength of the alloy, carbide is uniformly distributed in the alloy grain boundary, the grain boundary strength is improved, the dislocation slip speed is delayed, the purposes of improving the alloy strength, reducing the cracking tendency and the crack expansion rate are achieved, and the comprehensive performance of the alloy is ensured.
C) The nickel-based superalloy prepared by the preparation method provided by the invention has the following properties: room temperature performance: the tensile strength sigma b is more than or equal to 1300MPa (such as 1330 to 1360 MPa); the yield strength sigma 0.2 is equal to or more than 1050MPa (for example, 1052-1080 MPa); the elongation after break delta 5 is more than or equal to 27 percent (for example, 27 to 28 percent); the area reduction ratio psi is more than or equal to 35 percent (such as 35 to 40 percent); 730 ℃/550MPa durability performance: the lasting time tau is more than or equal to 70 hours (for example, 72 to 88 hours); the elongation after break delta 5 is more than or equal to 17.5 percent (for example, 17.7 to 20 percent); 815 ℃/295MPa durability: the lasting time tau is more than or equal to 83 hours (for example, 83 to 95 hours); the elongation after break delta 5 is more than or equal to 26 percent (for example, 26 to 30 percent); low cycle fatigue performance: 500 ℃ strain control 0-0.7%/0.33 Hz, > 5X 10 4 weeks (e.g., 57532-63645 weeks).
Additional features and advantages of the invention will be set forth in the description which follows, and in part will be obvious from the description, or may be learned by practice of the invention. The objectives and other advantages of the invention will be realized and attained by the structure particularly pointed out in the written description and claims thereof as well as the appended drawings.
Drawings
The drawings are only for purposes of illustrating particular embodiments and are not to be construed as limiting the invention, like reference numerals being used to refer to like parts throughout the several views.
FIG. 1 is a graph of crack growth rates for example 1 and comparative example 4;
FIG. 2 is a microstructure chart of example 1 of the present invention;
FIG. 3 is a second microstructure chart of example 1 of the present invention;
FIG. 4 is a microstructure of comparative example 4;
Fig. 5 is a grain diagram of comparative example 4.
Detailed Description
Preferred embodiments of the present invention are described in detail below with reference to the attached drawing figures, which form a part of the present invention and are used in conjunction with embodiments of the present invention to illustrate the principles of the present invention.
The invention provides a preparation method of a nickel-based superalloy, which comprises the following steps:
step 1: smelting to obtain an ingot;
Step 2: homogenizing heat treatment is carried out on the cast ingot; the homogenization heat treatment comprises 2 sections of heat preservation; 2, after the heat preservation of the 2 nd stage, furnace cooling to 1000+/-10 ℃ or lower, and then furnace discharging and air cooling;
Step 3: forging to prepare a bar blank and a forging;
Step 4: solution treatment;
step 5: stabilizing and treating with time effect to obtain nickel-base superalloy.
Specifically, the nickel-based superalloy comprises ：C：0.02％～0.06％,Cr：18.5％～20.0％,Co：13.0％～14.0％,Mo：4.0％～4.90％,Al：1.3％～1.6％,Ti：2.80％～3.25％,B：0.015％～0.03％,Zr：0.05％～0.15％,Ti/Al：2.25～2.38,(Al+Ti)：4.35％～4.58％,Fe：0.05％～2.0％,Mg≤0.005％,P：0.004％～0.009％,O：≤20PPm,N：≤20PPm,S≤10PPm, nickel in percentage by mass: the balance.
Specifically, in the step 1, the ingot is formed by adopting VIM+ESR+VAR triple smelting or VIM+VAR duplex smelting.
Specifically, in the step 2, the purpose of the homogenization heat treatment is to eliminate the low melting point phase and eutectic phase in the ingot and to reduce element segregation. The homogenization heat treatment is provided with 2 sections of heat preservation, and in the 1 st section of heat preservation, the temperature is lower, and the main effect is to eliminate a low-melting-point phase and a eutectic phase in the alloy; the heat preservation in the 2 nd stage can promote the uniform diffusion of segregation elements. Specifically, the homogenizing heat treatment process comprises the following steps:
S201, heating to 1165-1170 ℃, and preserving heat for 35-37 hours;
S202, continuously heating to 1200-1210 ℃, preserving heat for 38-42 h, cooling to below 1000+/-10 ℃, discharging and air cooling.
Specifically, in S201, the temperature is slowly raised, the furnace temperature is slowly raised to 1165-1170 ℃ from less than or equal to 400 ℃, and the temperature raising time is 10-15 h.
Specifically, in S201, considering that the low-melting-point phase and the eutectic phase formed by the B element in the as-cast structure are remelted due to the excessively high temperature and the excessively long heat preservation time, on one hand, holes are formed after homogenization to affect the subsequent alloy cogging and even cause ingot cracking to be scrapped, on the other hand, the B element cannot be fully diffused into the matrix to prepare for the subsequent formation of the grain boundary strengthening phase, and finally the grain boundary cannot be strengthened to improve the lasting creep property of the alloy. Therefore, the temperature is controlled to be raised to 1165-1170 ℃ and kept for 35-37 hours.
Specifically, in S202, the higher the temperature and the longer the heat preservation time, the more fully the remelting of the segregation element, but when the temperature and the time reach the matching balance point, the remelting of the segregation element will remain stable, so that the temperature is too high and the heat preservation time is too long, the positive effect on the remelting of the segregation element is smaller, and on the contrary, coarse grains, energy waste and production efficiency are caused; and the low temperature and the short heat preservation time can not ensure that most of segregation elements are remelted, and dendrite segregation reduces the hot working plasticity in the forging process. Therefore, the temperature is controlled to be raised to 1200-1210 ℃, and the temperature is kept for 38-42 h.
Specifically, in S202, considering that more grain boundary strengthening elements and carbide forming elements are added into the alloy of the invention, the strength of the alloy is improved, and meanwhile, the second phase is increased, and the risk of cracking of the alloy in the ingot casting and cogging process is increased, so that furnace cooling is controlled to 1000±10 ℃ and then furnace discharging and air cooling are performed, the furnace cooling treatment can enable the bent grain boundary to have more proportion and better high-temperature plasticity, so that the furnace cooling is controlled to be increased by controlling homogenization treatment, the toughness of the ingot casting in the cogging process is improved, and cracking is avoided.
Specifically, in the step 3, the forging method includes, but is not limited to, free forging or upsetting-drawing+radial forging combined cogging and subsequent hot rolling and cold drawing, wherein the fast forging method includes, but is not limited to, a combined method of single drawing and upsetting-drawing. The bar specification of cogging forging or subsequent rolling and cold drawing can be achievedRange.
Specifically, in the step 3, for the individual drawing, the drawing deformation amount is 30-60% per time, and the target size can be achieved by three to four times of drawing; for upsetting and drawing, the upsetting deformation is 30-60% per time, and the drawing deformation is 30-50% per time. The temperature range for preparing the bar is 1060-1200 ℃. The temperature of each fire is reduced by 40-80 ℃ until the temperature of the last fire is lower than 1080 ℃, wherein upsetting and drawing are carried out in a group, and then the temperature is reduced.
Specifically, in the step 3, the rod blank may be subjected to die forging to prepare a turbine disc with a disc shaft integrally or to prepare an annular member by ring rolling.
Specifically, in the step 4, the solution treatment process is as follows: heating to 995-1020 ℃, preserving heat for 3-5 h, discharging and cooling.
Specifically, in the step 4, the heat preservation time is insufficient because the alloy core does not reach the solid solution temperature due to the excessively high heating speed; therefore, the temperature is slowly raised, and the temperature raising speed is controlled to be 3-6 ℃/min. Too high a temperature or too long a holding time can result in grain growth that is detrimental to alloy performance. The temperature is low or the heat preservation time is too short, and the preparation for obtaining a proper gamma' phase in the subsequent aging process cannot be made; therefore, the temperature is controlled to be increased to 995-1020 ℃ and kept for 3-5 hours.
Specifically, in the step 4, in order to prevent the excessive stress caused by the too fast cooling speed from deforming the parts, different cooling modes are adopted for the parts with different wall thicknesses. For example, for disc shaft parts with the wall thickness of 20-200 mm, discharging into cooling oil for cooling; and (3) discharging air cooling is adopted for the annular piece with the wall thickness of 10-100 mm.
Specifically, in the step 5, the stabilizing and time-efficient treatment process comprises: heating to 820-850 ℃ for heat preservation for 5-7 h, air cooling to 500+/-10 ℃, continuously heating to 740-770 ℃ for heat preservation for 8-10 h, and air cooling to room temperature.
Specifically, in the step 5, considering that the temperature rising speed is too high, the alloy core does not reach the stabilization and aging treatment temperature, the heat preservation time is insufficient, the precipitated gamma' phase is insufficient, and the quantity and the size do not reach the optimal matching value; too high a temperature or too long a holding time can lead to larger gamma' -phase size, deviate from the optimal size-to-number ratio and affect the alloy strength. Therefore, the temperature is slowly increased, the temperature is increased from the temperature of the furnace to be less than or equal to 400 ℃, and the temperature increasing speed is controlled to be 2-3.5 ℃/min.
Specifically, in the step 5, the temperature is kept between 820 and 850 ℃ for 5 to 7 hours, and the temperature is cooled to 500+/-10 ℃ by air instead of being cooled to room temperature, so that the efficiency can be improved, and the size of carbide can be effectively ensured. For example, MC carbides having a size of 5 to 10 μm; the M 23C6 carbide has a size of about 0.3 to 1 μm.
Specifically, the O+N+S in the components of the nickel-based superalloy is less than or equal to 50PPm, for example, the O+N+S is less than or equal to 40PPm.
The following is a specific description of the action and the selection of the amount of the components contained in the nickel-based superalloy of the present invention:
c: the C content increases, and the carbide (MC, M 23C6) in the alloy also increases, but the precipitation temperature of M 23C6 does not change much. According to the invention, the beneficial phase M 23C6 is precipitated in the grain boundary by controlling the content of C and matching with the element B, so that the beneficial phase M 23C6 is continuously or necklace-shaped, film-shaped precipitation is avoided, and the film-shaped continuous precipitation is harmful to alloy performance, while the continuous or necklace-shaped precipitation is beneficial to alloy performance; by properly reducing the content of C and simultaneously controlling the content of alloy impurity elements, the contents of primary carbide and nitride in the crystal are effectively reduced, the source of fatigue cracking cracks is reduced, and the fatigue performance of the alloy is improved. When the content of C is generally more than 0.06%, the lasting life of the alloy is reduced, but the grain boundary segregation of B element is used for making the grain boundary membranous M 23C6 not easy to be separated out by improving the content of B element, zr element and Mg element, so that the grain boundary is improved. Therefore, the content of C is controlled to be 0.02-0.06%.
Cr: the Cr element can improve the oxidation resistance and corrosion resistance of the alloy. Meanwhile, cr is taken as a main forming element of M 23C6, the content of the Cr has small influence on the content of M 23C6, but the precipitation temperature of M 23C6 is obviously influenced, the more the content of Cr elements is, the higher the precipitation temperature of M 23C6 is, and when the content of Cr exceeds 20%, the precipitation temperature of M 23C6 exceeds 960 ℃. Therefore, the Cr content is controlled to be 18.5-20.0 percent in the invention.
Co: the Co element mainly plays a solid solution strengthening role, and the addition of Co to the gamma matrix can reduce the stacking fault energy of the matrix, reduce the stacking fault energy, increase the occurrence probability of the stacking fault and make the intersection and sliding of dislocation more difficult, so that the deformation requires larger external force and is shown as the improvement of the strength; and the stacking fault energy is reduced, the creep rate is reduced, and the creep resistance is increased. In addition, the Co element can also reduce the solubility of the gamma prime forming elements Ti and Al in the matrix, so that the quantity of gamma prime precipitated phases in the alloy is increased, and the service temperature of the alloy is increased. However, co belongs to a scarce resource in China, and the cost factor of the alloy is considered, so that the content of Co element is controlled to be 13.0-14.0%.
Mo: mo enters the nickel-based alloy matrix and mainly plays a role in solid solution strengthening. Meanwhile, the lattice constant of the matrix is increased, the gamma-gamma' mismatching degree is reduced, and the alloy stability and high-temperature creep property are improved. However, the increase in Mo element lowers the initial and final melting temperatures of the alloy. Therefore, the content of Mo element is controlled to be 4.0-4.9%.
Al: the Al element is the main forming element of the gamma' phase, and about 20% of Al added into the alloy enters a gamma solid solution to play a role in solid solution strengthening; 80% of Al and Ni form Ni 3 Al to play a role in precipitation strengthening. In order to ensure that the alloy has the necessary gamma prime phase to maintain the high temperature strength at 700 c. The invention controls the Al content to be 1.3-1.6%.
Ti: about 10% of the Ti added into the nickel-based alloy enters a gamma solid solution to play a certain solid solution strengthening role, and about 90% enters a gamma' phase. Under the condition of a certain Al content, as the quantity of the gamma 'phase is increased and the high-temperature strength of the alloy is increased along with the increase of the Ti content, in order to ensure that the gamma' phase which is necessary for maintaining the high-temperature strength of 700 ℃ exists in the alloy, the invention controls the Ti content to be 2.80-3.25 percent and simultaneously controls the Ti/Al:2.25 to 2.38, (Al+Ti): 4.35 to 4.58 percent.
Increasing the content of Al or Ti alone or simultaneously does not necessarily increase the content of the secondary γ' phase at the corresponding solid solution temperature. However, the high (Al+Ti) content has obvious advantages that the primary gamma 'phase is completely dissolved in the low (Al+Ti) group due to the improvement of the total dissolution temperature of the primary gamma' phase, the secondary gamma 'phase of the high (Al+Ti) group has more content in the temperature range where the primary gamma' phase is not completely dissolved in the high (Al+Ti) group, and the reasonable temperature control can prevent the grain size from growing due to the pinning effect of the primary gamma 'phase, and meanwhile, the gamma' phases of the two sizes can jointly play a reinforcing role.
B: boride (mainly comprising M 3B2,M4B3,MB12) is uniformly distributed in the crystal by homogenization treatment and cogging. The B content is increased, the segregation of B atoms in the grain boundary is increased, the binding force of the grain boundary is increased, meanwhile, the activation energy of grain boundary diffusion is increased, and the thermodynamic transformation of the B content is utilized to form beneficial continuous or necklace-shaped M 23C6 -type carbide in the grain boundary, so that the surface energy of grain boundary cracks can be increased while the grain boundary sliding is hindered, the grain boundary strength of the alloy is improved, the high-temperature durable creep performance is improved, the creep limit, particularly the durable strength of the material is obviously improved, and the higher the B content is, the better the high-temperature durable performance can be. However, when the B content is higher, the formation of a more multilayered low-melting phase is unfavorable for the improvement of the alloy properties but leads to the decrease of the properties, so that the B content is not excessively high. Comprehensively considering, the content of B is controlled to be 0.015-0.03 percent.
Fe: the lattice constant of Fe element is 3% different from that of Ni element, and its lattice expansion causes long-range stress field to block dislocation movement; meanwhile, fe can also reduce stacking fault energy, improve yield strength and play a solid solution strengthening role. The nickel-based superalloy of the invention has the elements mostly of grain boundary strengthening elements, and when the grain boundary strengthening elements are added, a part of Fe element is added, so that the strength of the substrate is improved and matched; thus, in the present invention, fe:0.05 to 2.0 percent.
P: according to the invention, the beneficial precipitated phase of the grain boundary P element is obtained by reasonably controlling the content of the P element, so that the high-temperature durable creep property of the alloy is improved, and meanwhile, the plasticity of the alloy is not damaged; but high levels of P have significant notch sensitivity for the high temperature durability of some alloys. In the invention, as a proper amount of B, zr elements are added, and comprehensive consideration is given, the control of P is needed in the invention: 0.004 to 0.009 percent.
Zr: the Zr element is used as a crystal boundary strengthening element, so that the solidus temperature of the nickel-based alloy can be effectively reduced, the solidification zone of the alloy can be enlarged, meanwhile, the C, S element can be promoted to be biased at the crystal boundary together with the Mg element, the alloy is purified, the creep resistance of the alloy can be improved by adding a proper amount of Zr element, but a large amount of eutectic phases such as Ni 5 Zr and the like can be formed when the content of the Zr element is too high, and the alloy is adversely affected by the alloy Jin Suxing. The invention improves the content of the element B and the content of the element Zr, and improves the strength of the material through the segregation of the Zr at the interface of the eutectic structure and the matrix and the grain boundary. For the problem of material plasticity reduction, the eutectic phase is eliminated by homogenization treatment regulation and control, and the eutectic phase is prevented from becoming a crack source. Considering comprehensively, zr needs to be controlled in the invention: 0.05 to 0.15 percent.
Mg: the Mg element has high affinity with S and is easy to form sulfide, so that the harmful risk of S is reduced; in addition, mg element is strongly segregated in the grain boundary and the phase boundary, and the carbide is thinned through the action of the grain boundary and the phase boundary, so that the grain boundary energy and the phase boundary energy are reduced, the interface binding force is improved, the creep strength and the plasticity of the alloy are improved, and the crack nucleation and the crack expansion are difficult to carry out. In the invention, the Mg content is controlled to be less than or equal to 0.005 percent by comprehensive consideration.
O: the reduction of O, N element content can reduce the quantity of inclusions in the material, and is beneficial to improving the plasticity and toughness of the alloy. During melting of the alloy, N readily forms Ti (C, N) with Ti, and the increase in Ti (C, N) increases the likelihood of fatigue source formation. Therefore, in the present invention, it is necessary to control O: less than or equal to 20PPm, N: and the volume of the mixture is less than or equal to 20PPm.
S: higher S content has an effect on the plasticity and long-term properties of the alloy. The higher the S content at the end of alloy smelting solidification, the easier the sulfide is to precipitate. The S element has obvious influence on the nickel-based alloy at the temperature of more than 800 ℃, the S element is obvious in the steel ingot smelting and cogging process, and the inventor finds that in the research: the 100ppm S experimental group fails to smelt, while the steel ingot of the group with the S content of 56ppm has serious cracking in the cogging process, and after the S content exceeds 10ppm, the durable service life and plasticity at 730 ℃ are reduced to different degrees. Therefore, in the present invention, it is necessary to control S: less than or equal to 10PPm.
In order to further improve the comprehensive performance of the nickel-base superalloy, the nickel-base superalloy comprises ：C：0.04％～0.06％,Cr：18.5％～20.0％,Co：13.0％～14.0％,Mo：4.0％～4.8％,Al：1.3％～1.5％,Ti：3.0％～3.25％,B：0.015％～0.03％,Zr：0.05％～0.15％,Ti/Al：2.3～2.38,(Al+Ti)：4.4％～4.58％,Fe：0.5％～1.0％,Mg≤0.005％,P：0.004％～0.009％,O：≤20PPm,N：≤20PPm,S≤10PPm, nickel in percentage by mass: the balance.
Specifically, the microstructure of the nickel-based superalloy mainly comprises equiaxed austenite grains, boride and carbide which are uniformly dispersed and distributed, and gamma' phase which is dispersed and distributed in the grains; wherein the carbide mainly comprises M 23C6 and MC; wherein the content of M 23C6 is 0.85-1.17%, the content of MC is 0.11-0.33%, and the content of boride is 0.35-0.5%; the size of the gamma 'phase is small, and the size of the gamma' phase is about 30-50 nm; the MC carbide has a size of about 5 to 10 μm and M 23C6 has a size of about 0.3 to 1. Mu.m.
Specifically, M 23C6 in the microstructure of the nickel-base superalloy is distributed in a continuous or necklace-like shape and a short rod-like shape, and is not precipitated in a film shape; MC is often in the form of a block with rounded boundaries.
Specifically, the nickel-based superalloy has uniform and fine crystal grains, and the average grain size can reach about 8-9 levels.
Specifically, the performance of the nickel-base superalloy obtained in the above step 5 is as follows: room temperature performance: the tensile strength sigma b is more than or equal to 1300MPa (such as 1330 to 1360 MPa); the yield strength sigma 0.2 is equal to or more than 1050MPa (for example, 1052-1080 MPa); the elongation after break delta 5 is more than or equal to 27 percent (for example, 27 to 28 percent); the area reduction ratio psi is more than or equal to 35 percent (such as 35 to 40 percent); 730 ℃/550MPa durability performance: the lasting time tau is more than or equal to 70 hours (for example, 72 to 88 hours); the elongation after break delta 5 is more than or equal to 17.5 percent (for example, 17.7 to 20 percent); 815 ℃/295MPa durability: the lasting time tau is more than or equal to 83 hours (for example, 83 to 95 hours); the elongation after break delta 5 is more than or equal to 26 percent (for example, 26 to 30 percent); low cycle fatigue performance: 500 ℃ strain control 0-0.7%/0.33 Hz, > 5X 10 4 weeks (e.g., 57532-63645 weeks).
The invention also provides application of the nickel-based superalloy, and the nickel-based superalloy can be used for rotating parts of aeroengines in service at the temperature of more than 700 ℃. Such as for alloy disk shaft integrated turbine disks.
Compared with the prior art, in the preparation method of the nickel-based superalloy, the technological parameters of each step are precisely controlled, for example, homogenization heat treatment comprises 2 sections of heat preservation; the cracking tendency is reduced by furnace discharging air cooling after the 2nd section is insulated and furnace cooling is carried out to below 1000+/-10 ℃; the combination of the accurate control of other steps and parameters of the steps avoids deformation, ensures that crystal grains of the alloy are uniform and tiny, ensures the best matching of the morphology and distribution of carbides in the alloy and the content and the size of other second phases, achieves the purposes of improving the strength of the alloy, reducing the cracking tendency and the crack expansion rate, and ensures the comprehensive performance of the nickel-based superalloy.
According to the preparation method of the nickel-based superalloy, the solid solution strengthening effect of the alloy and the grain boundary strength of the alloy are improved by accurately controlling the content of B, zr, fe, C, cr, co, al, ti and other single elements in the alloy; and by cooperatively controlling the values of B, zr and C, the best matching of the morphology and distribution of carbide in the alloy and the content and the size of other second phases is ensured, the risk crack sources in the alloy are reduced on the premise of not reducing the strength of the alloy, carbide is uniformly distributed in the alloy grain boundary, the grain boundary strength is improved, the dislocation slip speed is delayed, the purposes of improving the alloy strength, reducing the cracking tendency and the crack expansion rate are achieved, and the comprehensive performance of the alloy is ensured.
The nickel-based superalloy prepared by the preparation method provided by the invention has the following properties: room temperature performance: the tensile strength sigma b is more than or equal to 1300MPa (such as 1330 to 1360 MPa); the yield strength sigma 0.2 is equal to or more than 1050MPa (for example, 1052-1080 MPa); the elongation after break delta 5 is more than or equal to 27 percent (for example, 27 to 28 percent); the area reduction ratio psi is more than or equal to 35 percent (such as 35 to 40 percent); 730 ℃/550MPa durability performance: the lasting time tau is more than or equal to 70 hours (for example, 72 to 88 hours); the elongation after break delta 5 is more than or equal to 17.5 percent (for example, 17.7 to 20 percent); 815 ℃/295MPa durability: the lasting time tau is more than or equal to 83 hours (for example, 83 to 95 hours); the elongation after break delta 5 is more than or equal to 26 percent (for example, 26 to 30 percent); low cycle fatigue performance: 500 ℃ strain control 0-0.7%/0.33 Hz, > 5X 10 4 weeks (e.g., 57532-63645 weeks).
Examples 1 to 5
The advantages of the preparation method of the nickel-base superalloy of the present invention are shown below in specific examples and comparative examples. Embodiments 1-5 of the present invention provide a method for preparing a nickel-base superalloy.
The composition of the nickel-base superalloys of examples 1-5 is shown in Table 1 below.
The preparation method of the nickel-base superalloy of examples 1 to 5 includes:
Example 1:
Step 1: sequentially carrying out vacuum induction melting, electroslag remelting and vacuum arc remelting to obtain an ingot;
Step 2: homogenizing and annealing the cast ingot: heating from 400 ℃ to 1165 ℃ through 11h, and preserving heat for 37h; continuously heating to 1200 ℃, preserving heat for 38 hours, cooling to 1000+/-10 ℃, discharging and air cooling;
Step 3: preparing a bar blank by adopting quick forging and radial forging: sequentially upsetting and drawing out the alloy rod blank with the temperature of 1200 ℃ for three times, and radially forging; wherein, the deformation amount of upsetting each time is 30 percent, and the deformation amount of drawing each time is 30 percent; cooling to 40 ℃ after upsetting and drawing out each time, and cogging on a radial forging machine for multiple times until the heat preservation temperature is reduced to 1080 ℃ to obtain a finished bar with phi 180 mm; forging the bar blank by adopting a die forging press to prepare a forging piece: orderly upsetting cakes and die forging the alloy rod blank section with the temperature kept at 1060 ℃; the deformation amounts of the upsetting cakes and the die forging are respectively 50% and 35%, so that a forging piece is obtained;
step 4: solution treatment: heating from the furnace temperature of 400 ℃ at a heating speed of 5 ℃/min to 1015 ℃, preserving heat for 3.5h, and cooling with oil;
Step 5: stabilizing and performing time-efficient treatment to obtain a nickel-based superalloy disc member: heating from the furnace temperature of 400 ℃ at a heating speed of 5 ℃/min, heating to 850 ℃ and preserving heat for 5h, air-cooling to 500+/-10 ℃, continuously heating to 760 ℃ and preserving heat for 8h, and air-cooling to room temperature to obtain the nickel-based superalloy disc.
Example 2
The preparation method of this example is substantially the same as that of example 1, except that:
step 1: sequentially carrying out vacuum induction melting and vacuum arc remelting to obtain an ingot;
Step 3: upsetting, punching and ring rolling alloy rod blank sections with the temperature kept at 1060 ℃ in sequence; upsetting and ring rolling deformation are respectively 70% and 40%, so that a forging is obtained;
Step 4: cooling the forge piece after solid solution by adopting air cooling;
The remaining steps are the same.
Example 3
The preparation method of this example is substantially the same as that of example 1, except that:
Step 1: sequentially carrying out vacuum induction melting, electroslag remelting and vacuum arc remelting to obtain an ingot;
Step 3: preparing a bar blank by adopting a quick forging mode: sequentially upsetting and drawing out the alloy rod blank with the temperature of 1180 ℃ for three times, and radially forging; wherein, the deformation amount of upsetting each time is 50 percent, and the deformation amount of drawing each time is 50 percent; cooling to 40 ℃ after upsetting and drawing out each time until the heat preservation temperature is reduced to 1060 ℃ and then cogging to obtain a finished bar with phi 240 mm; forging the bar blank by adopting a die forging press to prepare a forging piece: orderly upsetting cakes and die forging the alloy rod blank section with the temperature kept at 1060 ℃; the deformation amounts of the upsetting cakes and the die forging are respectively 50% and 35%, so that a forging piece is obtained;
The remaining steps are the same.
Examples 4-5 differ from example 3 in the chemical composition of the alloy, see Table 1, with the remaining parameters being the same.
The invention also provides 6 comparative examples, examples 1-5 and comparative examples 1-5 steel with chemical compositions shown in Table 1.
Comparative examples 1 to 4 were prepared in the same manner as in example 1, except that the components were different.
Comparative example 5 has the same components as in example 1, and the homogenization heat treatment mode in the preparation method is different, wherein the homogenization method in comparative example 5 is to heat up to 1160 ℃ from 400 ℃ for 13 hours, and heat preservation is carried out for 48 hours; continuously heating to 1190 ℃, preserving heat for 66h, and then air cooling.
Comparative example 6 has the same components as example 1, and the homogenization heat treatment mode in the preparation method is different, wherein the homogenization method in comparative example 6 is that the homogenization is carried out after heat preservation for 38 hours, and then the homogenization is carried out in an air cooling furnace; cracks appear during the subsequent forging process.
TABLE 1 chemical composition wt%
The metallographic structures of the examples and comparative examples are shown in Table 2.
TABLE 2 metallographic structure
Table 3 shows the room temperature mechanical properties of the inventive examples and comparative examples, table 4 shows the durability properties of the inventive examples and comparative examples, and Table 5 shows the low cycle fatigue properties of the inventive examples and comparative examples.
TABLE 3 tensile mechanical Properties at room temperature
| Numbering device | σb/MPa | σ0.2/MPa | δ5/％ | ψ/％ |
| Example 1 | 1350 | 1080 | 27 | 39 |
| Example 2 | 1341 | 1062 | 28 | 37 |
| Example 3 | 1333 | 1052 | 27 | 35 |
| Example 4 | 1335 | 1054 | 28 | 40 |
| Example 5 | 1337 | 1061 | 28 | 38 |
| Comparative example 1 | 1411 | 1050 | 21 | 29 |
| Comparative example 2 | 1296 | 1018 | 23 | 32 |
| Comparative example 3 | 1231 | 1020 | 22 | 25 |
| Comparative example 4 | 1353 | 1010 | 24 | 36 |
| Comparative example 5 | 1206 | 1013 | 23 | 27 |
Table 4 durability properties
TABLE 5 500 ℃/strain control 0 to 0.7%/0.33Hz low cycle fatigue performance
| Numbering device | Cycle number Nf |
| Example 1 | 63645 |
| Example 2 | 60416 |
| Example 3 | 61013 |
| Example 4 | 62697 |
| Example 5 | 57532 |
| Comparative example 1 | 33884 |
| Comparative example 2 | 41348 |
| Comparative example 3 | 34438 |
| Comparative example 4 | 36428 |
| Comparative example 5 | 47323 |
TABLE 6 da/dN values at different ΔK (stress intensity factor)
FIG. 1 is a graph showing crack growth rates of example 1 and comparative example 4, and it can be seen that the da/dN value (fatigue crack growth rate) corresponding to the method is smaller under the same ΔK (stress intensity factor), and it is clear that the crack growth rate of the nickel-base superalloy prepared by the preparation method of the present invention is low.
Both fig. 2 and 3 are microstructure diagrams of example 1, in which the blocks MC with smooth boundaries are distributed in the crystal as seen in fig. 2, and in which the short bars M 23 C6 are continuously and uniformly distributed on the grain boundary as seen in fig. 3. Fig. 4 is a microstructure of comparative example 4 in which a film-like M 23C6 is distributed on the grain boundary, and fig. 5 is a grain diagram of comparative example 4, and it can be seen that the grain structure of comparative example 4 is a mixed crystal.
The present invention is not limited to the above-mentioned embodiments, and any changes or substitutions that can be easily understood by those skilled in the art within the technical scope of the present invention are intended to be included in the scope of the present invention.
Claims (10)
1. A method for preparing a nickel-based superalloy, comprising:
step 1: smelting to obtain an ingot;
Step 2: homogenizing heat treatment is carried out on the cast ingot; the homogenization heat treatment comprises 2 sections of heat preservation; 2, after the heat preservation of the 2 nd stage, furnace cooling to 1000+/-10 ℃ or lower, and then furnace discharging and air cooling;
Step 3: forging to prepare a bar blank and a forging;
Step 4: solution treatment;
step 5: stabilizing and treating with time effect to obtain nickel-base superalloy.
2. The method according to claim 1, wherein in the step 2, the homogenizing heat treatment comprises the steps of:
S201, slowly heating the furnace temperature to 1165-1170 ℃ from the temperature of less than or equal to 400 ℃ for 10-15 h, and preserving heat for 35-37 h;
S202, continuously heating to 1200-1210 ℃, preserving heat for 38-42 h, cooling to below 1000+/-10 ℃, discharging and air cooling.
3. The method according to claim 1, wherein in the step 4, the solution treatment process is as follows: heating to 995-1020 ℃, preserving heat for 3-5 h, discharging and cooling.
4. The method according to claim 3, wherein in the step 4, the temperature rising rate is 3 to 6 ℃/min.
5. The method according to claim 3, wherein in the step 4, the disc shaft parts with the wall thickness of 20-200 mm are cooled by discharging cooling oil; and (3) discharging air cooling is adopted for the annular piece with the wall thickness of 10-100 mm.
6. The method according to claim 1, wherein in the step 5, the stabilizing and time-efficient treatment process comprises: heating to 820-850 ℃ for heat preservation for 5-7 h, air cooling to 500+/-10 ℃, continuously heating to 740-770 ℃ for heat preservation for 8-10 h, and air cooling to room temperature.
7. The method according to claim 6, wherein in the step 5, the temperature rising rate is controlled to be 2 to 3.5 ℃/min.
8. The production method according to any one of claims 1 to 7, wherein the nickel-based superalloy comprises ：C：0.02％～0.06％,Cr：18.5％～20.0％,Co：13.0％～14.0％,Mo：4.0％～4.90％,Al：1.3％～1.6％,Ti：2.80％～3.25％,B：0.015％～0.03％,Zr：0.05％～0.15％,Ti/Al：2.25～2.38,(Al+Ti)：4.35％～4.58％,Fe：0.05％～2.0％,Mg≤0.005％,P：0.004％～0.009％,O：≤20PPm,N：≤20PPm,S≤10PPm, nickel in mass percent: the balance.
9. The method according to claim 8, wherein the microstructure of the nickel-base superalloy obtained in step 5 mainly comprises equiaxed austenite grains and uniformly dispersed boride and carbide, and gamma prime phase dispersed in the grains.
10. The method of claim 9, wherein the carbide is comprised mainly of M 23C6 and MC in the microstructure of the nickel-base superalloy.
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