A nickel-based heat-resistant alloy, its preparation method and application

Abstract

This invention belongs to the field of heat-resistant alloy technology, specifically relating to a nickel-based heat-resistant alloy, its preparation method, and its applications. The invention discloses a nickel-based heat-resistant alloy comprising: C: 0.05–0.08%, Cr: 21.0–23.0%, Co: 6.5–8.8%, Mo: 6.0–8.8%, Al: 0.70–1.40%, Ti: 0.40–0.80%, Nb: 0.15–0.55%, B: 0.002–0.005%, Zr: 0.015–0.06%, with the balance being nickel and unavoidable impurity elements, expressed as a percentage by mass. This alloy exhibits excellent strength, creep resistance, and structural stability, as well as good hot workability and weldability. It also has low raw material costs and is suitable for preparing ultra-supercritical turbine rotor forgings intended for long-term operation at 650°C.

Description

Technical Field

[0001] This invention belongs to the field of heat-resistant alloy technology, specifically relating to a nickel-based heat-resistant alloy, its preparation method, and its application. Background Technology

[0002] Currently, my country has begun tackling key technological challenges in high-parameter ultra-supercritical power plants, ranging from 630℃ to 650℃ and then to 700℃. The world's first 630℃ ultra-supercritical power plant demonstration project is under construction, and the technology for 650℃ ultra-supercritical power plants is undergoing feasibility studies and design. Heat-resistant materials are a major bottleneck restricting the development of thermal power units towards higher parameters. To meet the construction requirements of 650℃ ultra-supercritical power plants, it is necessary to develop heat-resistant materials for 650℃ ultra-supercritical steam turbine rotors.

[0003] The performance requirements of heat-resistant alloys for turbine rotor forgings in 650℃ ultra-supercritical thermal power units are mainly reflected in the following aspects: (1) Good high-temperature creep strength, 650℃ 10 5 (1) Extrapolated creep strength > 100 MPa; (2) Excellent high-temperature long-term structural stability; (3) Good room temperature and high-temperature instantaneous strength; (4) Good resistance to high-temperature steam oxidation; (5) Good cold and hot working properties; (6) Good welding properties. Studies have shown that martensitic heat-resistant steel can be used in ultra-supercritical steam turbine rotors at 630℃, but as the temperature continues to increase, martensitic heat-resistant steel cannot reach 650℃. 5 The requirement of an extrapolated creep strength greater than 100 MPa necessitates the use of nickel-based heat-resistant alloys for ultra-supercritical turbine rotor forgings operating at 650℃ and above. Therefore, it is necessary to research and develop nickel-based heat-resistant alloys to meet the requirements of ultra-supercritical turbine rotor forgings operating at 650℃ and above. Summary of the Invention

[0004] This invention is based on the inventor's discoveries and understanding of the following facts and problems:

[0005] The nickel-based heat-resistant alloy ingots required for 650℃ ultra-supercritical steam turbine rotor forgings weigh approximately 10-20 tons, with finished product diameters exceeding 1000 mm. However, nickel-based heat-resistant alloys have a narrow hot working temperature range, high deformation resistance, and are extremely difficult to manufacture. Therefore, it is necessary to develop and design nickel-based heat-resistant alloys for 650℃ ultra-supercritical steam turbine rotor forgings, taking into account both the service conditions of the 650℃ ultra-supercritical steam turbine rotor and the manufacturing process performance of the rotor forgings, as well as the mechanical properties of the heat-resistant materials.

[0006] This invention aims to at least partially solve one of the technical problems in related technologies. To this end, embodiments of this invention propose a nickel-based heat-resistant alloy that exhibits excellent strength, creep resistance, and structural stability, as well as good hot workability and weldability, and has low raw material costs. It is suitable for preparing ultra-supercritical turbine rotor forgings intended for long-term operation at 650°C.

[0007] The nickel-based heat-resistant alloy of this invention comprises: C: 0.05-0.08%, Cr: 21.0-23.0%, Co: 6.5-8.8%, Mo: 6.0-8.8%, Al: 0.70-1.40%, Ti: 0.40-0.80%, Nb: 0.15-0.55%, B: 0.002-0.005%, Zr: 0.015-0.06%, with the balance being nickel and unavoidable impurity elements, in mass percentage.

[0008] The advantages and technical effects of the nickel-based heat-resistant alloy in this invention are as follows: 1. In this invention, by controlling the content of Co and Mo elements, the segregation behavior in large ingots can be improved, the deformation resistance during forging can be reduced, the Mo element content can be controlled to avoid the formation of M6C, thus avoiding stress concentration and cracking caused by the formation of large-sized M6C during welding, and the Co element content can be controlled to reduce the coefficient of linear expansion of the alloy, thus avoiding cracking caused by the difference in thermal expansion coefficients during the welding of dissimilar materials, improving the weldability of the alloy and reducing production costs; 2. In this invention, the addition of Zr to the alloy can reduce grain boundary defects, improve grain boundary bonding, and form clustered ZrC around Ti(C,N), avoiding cracking caused by stress concentration and improving the durability of the alloy. However, when the amount of Zr added is low, it will preferentially enter the γ′ phase and cannot play a role in grain boundary strengthening and carbon formation. The role of compounds; therefore, in this embodiment of the invention, the addition amount of Zr is limited to 0.015-0.06%; 3. In this embodiment of the invention, Nb is added to the alloy to promote the formation of the MX phase and reduce the formation of M6C, thereby improving the weldability of the alloy. However, when the addition amount of Nb is too low, the improvement on the weldability of the alloy is very limited. Therefore, in this embodiment of the invention, the addition amount of Nb is limited to 0.15-0.55%; 4. In this embodiment of the invention, elements with strong segregation tendency such as W and V are not added to the alloy, which can improve the element segregation of large-tonnage alloy ingots, reduce the deformation resistance of nickel-based alloys, improve the weldability of alloys, and enable the alloy to have better ductility, toughness and creep performance while maintaining a high strength level; 5. In this embodiment of the invention, the alloy has good mechanical properties at both room temperature and high temperature. At room temperature of 20°C: tensile strength R m ≥870MPa, yield strength R p0.2 ≥480MPa, reduction of area A≥40%, elongation after fracture ≥40%, impact energy A kv≥110J; at a high temperature of 650℃: tensile strength R m ≥700MPa, yield strength R p0.2 ≥380MPa, reduction of area A≥41%, elongation after fracture≥42%, and extrapolated 100,000-hour creep strength at 650℃ is higher than 255MPa; 6. In the embodiments of the present invention, by controlling the amount of each element in the alloy within a suitable range, the strength of the alloy is maintained at a high level while having better plasticity and higher impact energy, which can be used to prepare ultra-supercritical steam turbine rotor forgings that have been in long-term service at 650℃.

[0009] In some embodiments, the impurity element includes at least one selected from Fe, Mn, Si, S, P, Cu, O, N, H, Pb, Sb, Sn, Bi, and As; the content of the impurity by mass percentage is: Fe≤0.5%, Mn≤0.10%, Si≤0.10%, S≤0.002%, P≤0.008%, Cu≤0.05%, N≤20ppm, O≤10ppm, H≤1ppm, Pb≤0.001%, Sb≤0.0025%, Sn≤0.0012%, Bi≤0.001%, and As≤0.002%.

[0010] In some embodiments, the contents of Co and Mo satisfy the following relationship: 6.7×Mo-50.7≤Co≤-1.65×Mo+25.3, where Co and Mo refer to the mass percentages of elements Co and Mo in the nickel-based heat-resistant alloy after removing the percentage sign.

[0011] In some embodiments, the contents of Al and Ti simultaneously satisfy the following relationships: 1.3≤Al+Ti≤1.7, 0.5≤Ti / Al≤0.8, where Al and Ti refer to the mass percentages of elements Al and Ti in the nickel-based heat-resistant alloy after removing the percentage sign.

[0012] In some embodiments, the contents of B and Zr satisfy the relationship: -15.6 < 3lgZr + 4lgB < -13.8, where B and Zr refer to the mass percentages of elements B and Zr in the nickel-based heat-resistant alloy after removing the percentage sign.

[0013] This invention also provides a method for preparing a nickel-based heat-resistant alloy, comprising the following steps:

[0014] (1) The raw materials are smelted, cast, homogenized and forged according to the design ratio to obtain a forging billet;

[0015] (2) The forging billet obtained in step (1) is subjected to solution treatment and aging treatment to obtain a nickel-based heat-resistant alloy.

[0016] The advantages and technical effects of the preparation method of nickel-based heat-resistant alloy in this invention are as follows: 1. The method of this invention, employing solution and aging heat treatment, combined with C and Nb elements in the alloy, can promote the growth of M. 23 The formation of C6 and MX phases, together with Al and Ti elements in the alloy, promotes the formation of γ′ phase, giving the alloy excellent room temperature and high temperature instantaneous strength; 2. The method of the present invention is simple, efficient and easy to operate, and the alloy obtained has excellent comprehensive performance, which can be used for the preparation of large-tonnage alloy ingots and is suitable for promotion and application in industrial production.

[0017] In some embodiments, in step (1), the smelting is carried out using a process of vacuum induction melting + vacuum consumable remelting or vacuum induction melting + protective atmosphere electroslag melting + vacuum consumable remelting.

[0018] In some embodiments, in step (1), the homogenization process is performed using a graded heating process.

[0019] In some embodiments, in step (2), the solution treatment includes heating the forging billet to 1150-1190°C and holding it for 60-90 minutes, followed by water cooling to room temperature; the aging treatment includes heating the forging to 700-750°C and holding it for 16-24 hours, followed by air cooling to room temperature.

[0020] This invention also provides the application of the above-mentioned nickel-based heat-resistant alloy or the nickel-based heat-resistant alloy prepared by the above-mentioned method in ultra-supercritical steam turbine rotors. Attached Figure Description

[0021] Figure 1 This is a schematic diagram showing the effect of Co and Mo content on the M6C and μ phases in nickel-based heat-resistant alloys;

[0022] Figure 2 This is a graph showing the relationship between the Ti / Al ratio and the degree of mismatch.

[0023] Figure 3 These are SEM images of carbonitrides in the alloy prepared in Example 1, where (a) is ZrC and (b) is clustered ZrC formed around Ti(C,N);

[0024] Figure 4 These are schematic diagrams of the sampling locations for Application Example 1 and Application Comparative Example 1;

[0025] Figure 5 This is a SEM image of the nickel-based heat-resistant alloy prepared in Example 1;

[0026] Figure 6 This is a SEM image of the nickel-based heat-resistant alloy prepared in Comparative Example 1.

[0027] Figure 7 This is a SEM image of the nickel-based heat-resistant alloy prepared in Comparative Example 2. Detailed Implementation

[0028] Embodiments of the present invention are described in detail below, examples of which are illustrated in the accompanying drawings. The embodiments described below with reference to the accompanying drawings are exemplary and intended to explain the present invention, and should not be construed as limiting the present invention.

[0029] The nickel-based heat-resistant alloy of this invention comprises: C: 0.05-0.08%, Cr: 21.0-23.0%, Co: 6.5-8.8%, Mo: 6.0-8.8%, Al: 0.70-1.40%, Ti: 0.40-0.80%, Nb: 0.15-0.55%, B: 0.002-0.005%, Zr: 0.015-0.06%, with the balance being nickel and unavoidable impurity elements, in mass percentage.

[0030] The nickel-based heat-resistant alloy of this invention improves segregation behavior in large ingots and reduces deformation resistance during forging by controlling the content of Co and Mo. Controlling the Mo content prevents the formation of M6C, avoiding stress concentration and cracking caused by large-sized M6C during welding. Controlling the Co content reduces the alloy's coefficient of linear expansion, preventing cracks caused by differences in thermal expansion coefficients during welding of dissimilar materials, thus improving weldability and reducing production costs. Adding Zr to the alloy reduces grain boundary defects, improves grain boundary bonding, and forms clustered ZrC around Ti(C,N), preventing cracking caused by stress concentration and improving the alloy's creep resistance. However, when the Zr content is low, it preferentially enters the γ′ phase, failing to provide grain boundary strengthening and carbide formation. Therefore, in this embodiment of the invention, the amount of Zr added is limited to 0.015-0.06%. Adding Nb to the alloy promotes the formation of the MX phase, reduces the formation of M6C, and improves the alloy's weldability. However, when the amount of Nb added is too low, the improvement in the alloy's weldability is very limited. Therefore, in this embodiment of the invention, the amount of Nb added is limited to 0.15-0.55%. Not adding elements with strong segregation tendency, such as W and V, to the alloy can improve elemental segregation in large-tonnage alloy ingots, reduce the deformation resistance of nickel-based alloys, improve the alloy's weldability, and, while maintaining a high strength level, give the alloy better ductility, toughness, and creep resistance. This alloy exhibits good mechanical properties at both room temperature and high temperature. At room temperature (20°C): tensile strength R... m ≥870MPa, yield strength R p0.2 ≥480MPa, reduction of area A≥40%, elongation after fracture ≥40%, impact energy A kv ≥110J; at a high temperature of 650℃: tensile strength R m≥700MPa, yield strength R p0.2 With a strength of ≥380MPa, a reduction of area A≥41%, an elongation after fracture≥42%, and an extrapolated creep strength of 100,000 hours at 650℃ higher than 255MPa, the alloy maintains a high strength while possessing better plasticity and higher impact energy by controlling the amount of each element in the alloy within a suitable range. This makes it suitable for ultra-supercritical turbine rotor forgings that can be used for long-term service at 650℃.

[0031] The roles of each element in the alloy of this invention are as follows:

[0032] Carbon (C): In nickel-based heat-resistant alloys, C mainly forms carbides. During aging treatment, the granular, discontinuous carbides precipitated at grain boundaries can prevent grain boundary sliding and crack propagation, thus improving the alloy's creep rupture life. Simultaneously, carbides significantly affect the alloy's instantaneous strength at both room temperature and high temperature. Rotors require high material strength, necessitating a certain carbon content to improve instantaneous strength at both temperatures. However, when C exceeds a certain threshold, the contribution of carbides to instantaneous strength decreases significantly. Furthermore, improper heat treatment processes or prolonged aging treatments can lead to a large amount of carbide precipitation due to microstructural instability, forming a continuous network along grain boundaries, which significantly reduces the material's ductility and toughness. Therefore, in this invention, the mass percentage of C is controlled within the range of 0.05% to 0.08%, specifically, for example, 0.05%, 0.06%, 0.07%, and 0.08%, to ensure that the carbide mass fraction is above 1%.

[0033] Aluminum (Al) and Titanium (Ti): By adding Al and Ti elements, a γ′(Ni3(Al,Ti)) phase is formed. The γ′ phase is dispersed in the matrix, hindering dislocation movement and significantly improving the material strength. In this invention, the mass percentage range of aluminum and titanium is controlled as follows: Al: 0.70–1.40%, specifically, for example, 0.70%, 0.80%, 0.90%, 1.00%, 1.10%, 1.20%, 1.30%, 1.40%; Ti: 0.40–0.80%, specifically, for example, 0.40%, 0.50%, 0.60%, 0.70%, 0.80%.

[0034] Niobium (Nb): Nb is a strong carbide-forming element, capable of forming stable NbC or Nb(C,N) phases, significantly improving the room temperature and high temperature instantaneous strength of materials. Furthermore, the MX phase has a high melting point and exhibits better stability than M during long-term service. 23C6; Elemental segregation occurs during welding. Mo and C are both interdendritic segregating elements, which may lead to the presence of M6C and rapid coarsening, causing stress concentration and resulting in welding cracks. Adding Nb can preferentially combine with carbon, reducing the formation of large-sized M6C phases, reducing stress concentration, avoiding welding crack formation, and improving welding performance. If the amount of Nb added is too small, there is no obvious MC phase precipitation during welding. If the amount of Nb added is too large, it will lead to grain boundary widening, forming a liquid film, and a large number of chain-like precipitates distributed at the grain boundaries, which easily leads to secondary welding cracks. Therefore, in this invention, the mass percentage of Nb is controlled in the range of 0.15% to 0.55%, specifically, for example, 0.15%, 0.20%, 0.30%, 0.40%, 0.50%, and 0.55%.

[0035] Chromium (Cr): The main functions of Cr are anti-oxidation and anti-hot corrosion; in addition, Cr is also a metal element. 23 The C6 phase is a major forming element, and the resulting carbides can hinder grain boundary movement, thereby improving the high-temperature creep strength of the material. However, excessively high Cr content in the alloy can promote the precipitation of the harmful σ phase. Therefore, in this invention, the mass percentage of Cr is controlled within the range of 21.0% to 23.0% to ensure sufficient oxidation resistance of the alloy while avoiding the formation of the σ phase. Specifically, for example, 21.0%, 21.5%, 22.0%, 22.5%, and 23.0%.

[0036] Molybdenum (Mo): Mo has a large atomic radius, 9-12% larger than that of Ni, Co, and Fe, and exhibits strong solid solution strengthening. However, excessive Mo addition leads to greater deformation resistance during alloy deformation. Therefore, controlling the amount of Mo added reduces this deformation resistance during forging, improving the alloy's hot working properties and facilitating the manufacture of large forgings. While Mo provides significant strengthening, excessive addition can lead to the precipitation of the harmful μ phase. Furthermore, Mo is prone to segregation, which is more severe in large-tonnage alloy ingots, resulting in inhomogeneous ingot microstructure. Mo's segregation coefficient K is less than 1, causing it to easily agglomerate between dendrites during solidification, combining with segregated C to form M6C carbides. M6C carbide aggregation can cause excessively high local residual stress, leading to weld solidification cracks and deteriorating weldability. Therefore, in this invention, the mass percentage of Mo is controlled within the range of 6.0% to 8.8%, specifically, for example, 6.0%, 6.5%, 7.0%, 7.5%, 8.0%, 8.5%, and 8.8%.

[0037] Cobalt (Co): Controlling the Co content can reduce the coefficient of linear expansion of the alloy, avoid cracks caused by differences in thermal expansion coefficients during welding of dissimilar materials, and improve the weldability of the alloy. Furthermore, Co can significantly increase the solid solubility of Mo in the nickel matrix. In this invention, the mass percentage of Co is controlled within the range of 6.5% to 8.8%, specifically, for example, 6.5%, 7.0%, 7.5%, 8.0%, 8.5%, and 8.8%.

[0038] Tungsten (W): Similar to Mo, W is a solid solution strengthening element with a better strengthening effect than Mo. However, W has a stronger tendency to segregate. Since the alloy ingots used in ultra-supercritical turbine rotors are in the 10-20 ton range, the introduction of W can easily lead to severe elemental segregation, impairing alloy properties. Furthermore, the alloy does not require excessively high levels of solid solution strengthening elements. Therefore, W is not added in this invention.

[0039] Boron (B): B is a grain boundary strengthening element that can improve the high-temperature creep strength of alloys. B will [strengthen the grain boundary] at γ′ / M [grain boundary]. 23 C6 and γ / M 23 In the C6 interface, it leans towards M. 23 The C6 phase is distributed on one side, which can reduce C segregation at grain boundaries, and some B can enter the carbide to form M. 23 (C,B)6, reduces M at grain boundaries 23 The coarsening rate of C6; too low a B content results in an insignificant strengthening effect, while too high a B content leads to the formation of grain boundary liquidation cracks in the low-melting-point boride phases (such as M3B2) generated during alloy welding solidification, resulting in deterioration of hot workability and weldability. Therefore, in this invention, the mass percentage of B is controlled within the range of 0.002–0.005%, specifically, for example, 0.002%, 0.003%, 0.004%, and 0.005%.

[0040] Zirconium (Zr): Zr can improve the structure and morphology of grain boundaries and strengthen them. The combined addition of Zr and boron helps to further enhance the grain boundary strengthening effect. The addition of Zr forms clusters of ZrC, which have a pinning effect on grain boundaries. ZrC forms around Ti(C,N). Ti(C,N) is mostly rectangular, and there is a large stress concentration at its tip with the matrix, which is prone to becoming a crack initiation site. ZrC precipitates in clusters around it, avoiding cracking caused by stress concentration. Zr can enter the γ′ phase, increasing the amount of γ′ phase and raising the γ′ phase dissolution temperature. If the Zr content is too low, most of Zr will enter the γ′ phase, and the grain boundary strengthening effect will not be obvious. If the Zr content is too high, it will lead to the formation of Zr- and Ni-rich intermetallic compound phases, which will cause the alloy to crack during hot working. Zr readily combines with oxygen to form Zr-rich oxide inclusions. Therefore, in this invention, the Zr element content ranges from 0.015% to 0.006%, specifically, for example, 0.015%, 0.02%, 0.03%, 0.04%, 0.05%, and 0.06%.

[0041] Magnesium (Mg): Mg can improve the creep rupture life and ductility of nickel-based alloys, but controlling the amount of Mg during the smelting process is difficult. Excessive Mg can easily form a low-melting-point Ni-Mg phase, leading to a decrease in hot working properties. Therefore, Mg is not added in this invention.

[0042] In this embodiment of the invention, unavoidable impurity elements and gaseous elements are also included, such as residues generated during the manufacturing process from raw materials, slag, and refractory materials, such as Fe, Mn, Si, S, P, Cu, O, N, H, Pb, Sb, Sn, Bi, and As. These elements need to be strictly controlled. By mass percentage, the controlled impurity content in the alloy is as follows: Fe≤0.5%, Mn≤0.10%, Si≤0.10%, S≤0.002%, P≤0.008%, Cu≤0.05%, N≤20ppm, O≤10ppm, H≤1ppm, Pb≤0.001%, Sb≤0.0025%, Sn≤0.0012%, Bi≤0.001%, and As≤0.002%.

[0043] In some embodiments, preferably, the contents of Co and Mo satisfy the following relationship: 6.7×Mo-50.7≤Co≤-1.65×Mo+25.3, where Co and Mo refer to the mass percentages of elements Co and Mo in the nickel-based heat-resistant alloy after removing the percentage sign.

[0044] In this embodiment of the invention, the solid solubility of Co and Mo elements in nickel-based alloys exhibits a matching relationship, such as... Figure 1As shown, when the Mo content is too high, if the Co content is below the critical value, it will lead to the formation of M6C carbides, which will deteriorate the welding performance and creep performance. If the Co content is above the critical value, it will lead to the formation of μ phase, which will reduce the strength of the alloy and cause the alloy to crack. In this embodiment of the invention, by optimizing the ratio of elements Co and Mo to satisfy the relationship 6.7×Mo-50.7≤Co≤-1.65×Mo+25.3, the formation of M6C carbides and μ phase can be avoided, and the microstructure stability of the material during long-term service can be improved.

[0045] In some embodiments, preferably, the contents of Al and Ti simultaneously satisfy the following relationships: 1.3≤Al+Ti≤1.7, 0.5≤Ti / Al≤0.8, where Al and Ti refer to the mass percentages of elements Al and Ti in the nickel-based heat-resistant alloy after removing the percentage sign.

[0046] In nickel-based heat-resistant alloys, the amount of γ′ phase increases with increasing Ti+Al content. To ensure alloy strength, the mass fraction of γ′ phase is required to be above 5%. The microstructure stability of the γ′ phase significantly affects alloy performance, and the stability of the γ′ phase is related to the absolute value of mismatch. In this invention, when the Al+Ti content is between 1.3% and 1.7% and the Ti / Al ratio is controlled between 0.5 and 0.8, the absolute value of mismatch is ≤0.03%. The alloy has a low mismatch, and the γ′ phase is relatively stable during high-temperature service, thus ensuring good microstructure stability of the alloy during long-term high-temperature service.

[0047] The method for calculating the mismatch degree is as follows:

[0048] A 10mm × 10mm × 5mm sample was cut, polished, and then electrolyzed with 10% perchloric acid alcohol to remove surface stress. XRD analysis was performed to determine the austenite (γ) lattice constant. The sample was then electrolyzed with LiCl solution to separate the γ′ phase, and XRD analysis was used to determine the γ′ phase lattice constant. The misfit degree was calculated according to the formula shown in equation (1):

[0049]

[0050] In the formula a γ′ and a γ are the lattice constants of the γ′ phase and the γ phase, respectively.

[0051] Mismatch degree affects the interfacial energy of the precipitated phases. When the absolute value of mismatch degree is low, the anisotropy of the interfacial energy is low, and the precipitated phase remains spherical throughout long-term service. The spherical γ′ phase has low energy and a slow coarsening rate, resulting in stable microstructure and properties during long-term service. As the absolute value of mismatch degree increases, the γ′ phase transforms into a cubic phase. Increased mismatch degree leads to increased lattice distortion, which, while beneficial to increasing alloy strength, causes the cubic γ′ phase to coarsen easily during service, leading to a rapid decline in its performance after service.

[0052] Depend on Figure 2 It is known that Ti / Al affects the mismatch degree of the alloy. In order to ensure the stability of the γ′ phase structure during long-term service, the absolute value of the mismatch degree should be controlled to keep it at a low level. When the Ti / Al ratio is between 0.5 and 0.8, the absolute value of the mismatch degree is ≤0.03%.

[0053] In some embodiments, preferably, the contents of B and Zr satisfy the relationship: -15.6 < 3lgZr + 4lgB < -13.8, where B and Zr refer to the mass percentages of elements B and Zr in the nickel-based heat-resistant alloy after removing the percentage sign.

[0054] In this embodiment of the invention, B and Zr are further restricted to satisfy the relationship -15.6 < 3lgZr + 4lgB < -13.8, which can give full play to the synergistic effect between B and Zr, further improve the grain boundary strengthening effect, and thus improve the performance of the alloy.

[0055] This invention also provides a method for preparing the above-mentioned nickel-based heat-resistant alloy, comprising the following steps:

[0056] (1) The raw materials are smelted, cast, homogenized and forged according to the design ratio to obtain a forging billet;

[0057] (2) The forging billet obtained in step (1) is subjected to solution treatment and aging treatment to obtain a nickel-based heat-resistant alloy.

[0058] The method for preparing the nickel-based heat-resistant alloy according to the embodiments of the present invention employs solution and aging heat treatment, and, in conjunction with the C and Nb elements in the alloy, can promote the growth of M. 23 The formation of C6 and MX phases, together with Al and Ti elements in the alloy, promotes the formation of γ′ phase, giving the alloy excellent room temperature and high temperature instantaneous strength. The process is simple, efficient and easy to operate, and the resulting alloy has excellent comprehensive properties. It can be used for the preparation of large-tonnage alloy ingots and is suitable for widespread application in industrial production.

[0059] In some embodiments, preferably, the smelting is carried out using a process of vacuum induction melting + vacuum consumable remelting or vacuum induction melting + protective atmosphere electroslag remelting + vacuum consumable remelting.

[0060] In this embodiment of the invention, a preferred smelting process is selected, which enables the alloy ingot to have high purity and uniform and dense structure, thereby obtaining a nickel-based heat-resistant alloy with better performance.

[0061] In some embodiments, preferably, the alloy ingot cast after smelting is not less than 10 tons, and more preferably 10-20 tons.

[0062] In this embodiment of the invention, by controlling the composition of the alloy, the segregation in the preparation of large-tonnage alloy ingots can be significantly reduced, enabling the invention to manufacture alloy ingots of 10 tons or more for use in the preparation of ultra-supercritical steam turbine rotors.

[0063] In some embodiments, preferably, the homogenization treatment is performed using a staged heating process. More preferably, the staged heating process includes: loading the furnace at a temperature below 700°C and homogenizing it at a rate of 0.5–1 mm / min; then raising the temperature to 990–1010°C at a rate of 30–60°C / h and holding it for 1–2 hours; then raising the temperature to 1120–1140°C at a rate of 25–50°C / h and holding it for 1–2 hours; finally raising the temperature to 1180–1200°C at a rate of 10–20°C / h to eliminate microsegregation; and then cooling the alloy ingot to below 600°C by furnace cooling after high-temperature homogenization.

[0064] In this embodiment of the invention, preferably, a graded heating process can be adopted based on the specific grain growth kinetics and thermal properties of nickel-based heat-resistant alloys, such as thermal conductivity, to avoid the problem of alloy ingot cracking caused by thermal stress due to excessive temperature difference.

[0065] In some embodiments, in step (2), the solution treatment includes heating the forging billet to 1150-1190°C and holding it at that temperature for 60-90 minutes, followed by water cooling to room temperature; the aging treatment includes heating the forging billet to 700-750°C and holding it at that temperature for 16-24 hours, followed by air cooling to room temperature.

[0066] The heat treatment method of this invention can control the grain size of the alloy to level 1-3.

[0067] This invention also provides the application of the above-mentioned nickel-based heat-resistant alloy or the nickel-based heat-resistant alloy prepared by the above-mentioned method in ultra-supercritical steam turbine rotors.

[0068] The technical solution of the present invention will be described in detail below with reference to the embodiments and accompanying drawings.

[0069] Example 1

[0070] The nickel-based heat-resistant alloy of this embodiment comprises: C: 0.065%; Cr: 21.92%; Co: 7.99%; Mo: 6.98%; Al: 1.00%; Ti: 0.64%; Nb: 0.29%; B: 0.0037%; Zr: 0.024%; N: 0.0014%; Si: 0.032%; Mn: 0.0022%; P: 0.0038%; S: 0.0013%; Cu: 0.0026%; Fe: 0.054%; H: <0.00005%; O: 0.0008%, with the balance being Ni, by mass percentage.

[0071] (1) According to the alloy composition ratio, the raw materials are smelted by vacuum induction melting + vacuum self-consumption remelting process and cast into 200Kg alloy ingots, and the content of impurity elements is strictly controlled.

[0072] (2) The furnace is loaded with heat at a temperature below 700℃ and homogenized at a rate of 0.6-0.8 mm / min. Then, the temperature is increased to 1000℃ at a rate of 40℃ / h and held for 1.5h. Then, the temperature is increased to 1130℃ at a rate of 35℃ / h and held for 1h. Finally, the temperature is increased to 1190℃ at a rate of 15℃ / h and held for 20h to perform high-temperature homogenization treatment and eliminate microsegregation. Then, the furnace is cooled to 600℃ and then air-cooled.

[0073] (3) The alloy ingot obtained in step 2 is forged into a billet, and then forged into a round bar with a diameter of 16 mm. The initial forging temperature is 1170℃, and the final forging temperature is 1050℃. If the temperature is lower than the final forging temperature during the forging process, it needs to be reheated in the furnace before forging again.

[0074] (4) The forged bar obtained in step 3 is subjected to heat treatment. The solution treatment is carried out at 1160℃ for 1 hour and then water-cooled to room temperature. Subsequently, it is carried out at 730℃ for 16 hours and then air-cooled to obtain a nickel-based heat-resistant alloy.

[0075] The alloy composition obtained in Example 1 is shown in Table 1, and the relationship between the elements in the alloy is shown in Table 2.

[0076] Examples 2-10

[0077] The preparation methods of Examples 2 to 10 are the same as those of Example 1, except that the alloy composition is different. The alloy composition obtained in Examples 2 to 10 is shown in Table 1, and the relationship between the elements in the alloy is shown in Table 2.

[0078] Comparative Examples 1-10

[0079] The preparation methods of Comparative Examples 1 to 10 are the same as those of Example 1, except that the alloy composition is different. The alloy composition of Comparative Examples 1 to 10 is shown in Table 1, and the relationship between the elements in the alloy is shown in Table 2.

[0080] Application Example 1

[0081] The chemical composition of this application embodiment, by mass percentage, includes: C: 0.064%; Cr: 21.94%; Co: 8.05%; Mo: 7.01%; Al: 0.97%; Ti: 0.64%; Nb: 0.29%; B: 0.0035%; Zr: 0.025%; N: 0.0015%; Si: 0.029%; Mn: 0.0021%; P: 0.0035%; S: 0.0013%; Cu: 0.0041%; Fe: 0.098%; H: <0.00005%; O: 0.0009%, with the balance being Ni.

[0082] (1) According to the alloy composition ratio, the alloy ingots are smelted by a three-stage process of vacuum induction melting, protective atmosphere electroslag remelting, and vacuum self-consumption remelting, and then cast into 10t alloy ingots with strict control of impurity element content.

[0083] (2) The furnace is loaded with heat at a temperature below 700℃ and homogenized at a rate of 0.6-0.8 mm / min. Then, the temperature is increased to 1000℃ at a rate of 40℃ / h and held for 1.5h. Then, the temperature is increased to 1130℃ at a rate of 35℃ / h and held for 1h. Finally, the temperature is increased to 1190℃ at a rate of 15℃ / h and held for 20h to perform high-temperature homogenization treatment and eliminate microsegregation. Then, the furnace is cooled to 600℃ and then air-cooled.

[0084] (3) The alloy ingot obtained in step 2 is forged into a billet and then forged into a columnar forging with a diameter of Φ900mm. The forging temperature is 1170℃ and the final forging temperature is 1050℃. If the temperature is lower than the final forging temperature during the forging process, it needs to be reheated in the furnace before forging.

[0085] (4) The forging obtained in step 3 is subjected to heat treatment. The solution treatment is held at 1160℃ for 1 hour and then water-cooled to room temperature. Subsequently, it is held at 730℃ for 16 hours and then air-cooled to obtain nickel-based heat-resistant alloy forging products.

[0086] Application Comparative Example 1

[0087] The preparation method of Comparative Example 1 is the same as that of Application Example 1, except that the alloy composition is different. The alloy composition of Comparative Example 1 is similar to that of Comparative Example 8.

[0088] The chemical composition of Comparative Example 1, by mass percentage, includes: C: 0.058%; Cr: 21.79%; Co: 11.93%; Mo: 8.77%; Al: 1.13%; Ti: 0.42%; Nb: 0.28%; W: 0.67%; B: 0.0030%; Zr: 0.026%; N: 0.0014%; Si: 0.031%; Mn: 0.0026%; P: 0.0039%; S: 0.0014%; Cu: 0.0047%; Fe: 0.081%; H: <0.00005%; O: 0.0008%, with the balance being Ni and unavoidable impurity elements.

[0089] Table 1 Chemical composition (%) of Examples 1-10 and Comparative Examples 1-10

[0090] alloy C Cr Co Mo Al Ti Nb V W B Zr Example 1 0.065 21.94 7.99 6.98 1.00 0.64 0.29 0 0 0.0037 0.024 Example 2 0.058 21.92 8.02 7.42 0.92 0.51 0.30 0 0 0.0033 0.046 Example 3 0.064 22.07 8.03 7.02 0.91 0.70 0.29 0 0 0.0028 0.023 Example 4 0.066 22.09 8.80 6.99 1.01 0.62 0.28 0 0 0.0049 0.022 Example 5 0.060 21.99 8.04 8.49 0.97 0.62 0.26 0 0 0.0031 0.024 Example 6 0.059 22.04 8.11 6.97 1.04 0.61 0.55 0 0 0.0029 0.023 Example 7 0.064 21.96 8.03 7.01 1.01 0.42 0.31 0 0 0.0033 0.021 Example 8 0.065 21.94 7.98 6.99 0.76 0.75 0.31 0 0 0.0031 0.02 Example 9 0.062 22.03 7.96 7.04 0.81 0.42 0.28 0 0 0.0031 0.021 Example 10 0.064 21.97 7.93 6.98 0.95 0.67 0.33 0 0 0.0049 0.045 Comparative Example 1 0.058 21.92 8.71 8.98 0.97 0.61 0.29 0 0 0.0033 0.024 Comparative Example 2 0.060 21.96 11.04 8.54 1.10 0.41 0.28 0 0 0.0032 0.022 Comparative Example 3 0.061 22.02 8.37 8.74 1.11 0.42 0.27 0.2 0 0.003 0.022 Comparative Example 4 0.058 21.98 8.78 8.78 0.61 0.54 0.31 0 0 0.0029 0.021 Comparative Example 5 0.041 22.09 8.49 8.75 0.92 0.62 0.29 0 0 0.0026 0.022 Comparative Example 6 0.092 21.88 8.81 8.69 0.91 0.65 0.21 0 0 0.0029 0.045 Comparative Example 7 0.061 22.03 8.09 8.67 0.90 0.64 0.22 0 0.41 0.0030 0.022 Comparative Example 8 0.058 21.94 11.87 8.74 1.13 0.40 0.19 0 0.4 0.0031 0.023 Comparative Example 9 0.064 21.97 8.04 7.02 1.02 0.61 0.29 0 0 0.0036 0 Comparative Example 10 0.050 21.59 8.04 7.14 1.24 1.28 0.20 0 0.8 0.0024 0.022

[0091] Table 2. Relationships of alloying elements in Examples 1-10 and Comparative Examples 1-10

[0092]

[0093]

[0094] Performance testing

[0095] (1) The room temperature tensile properties, high temperature tensile properties, room temperature impact properties, 650℃ high temperature creep properties and weldability of the alloys prepared in Examples 1-10 and Comparative Examples 1-10 were tested, and the results are shown in Tables 3-7.

[0096] The welding performance was tested using the same specification welding wire and low-energy-inert-gas tungsten inert gas with a narrow bevel. Liquefaction cracks and stress relaxation cracks in the heat-affected zone (HAZ) of each weld joint were observed. Longitudinal samples from the welded surfaces were ground, polished, etched, and then examined using a metallographic microscope.

[0097] Table 3. Room temperature tensile properties (20°C) of Examples 1-10 and Comparative Examples 1-10

[0098]

[0099]

[0100] Table 4. High-temperature tensile properties (650℃) of Examples 1-10 and Comparative Examples 1-10

[0101] alloy Tensile strength / MPa Yield strength / MPa Elongation after fracture / % Reduction of area / % Example 1 840 440 46 47 Example 2 805 445 46.5 46.5 Example 3 840 470 44 42.5 Example 4 835 430 45 48 Example 5 825 435 44 45 Example 6 820 430 46 46.5 Example 7 845 460 42 41.5 Example 8 855 475 42 41 Example 9 715 380 50 49.5 Example 10 845 440 43 45 Comparative Example 1 815 405 50 51.5 Comparative Example 2 845 435 45 47 Comparative Example 3 840 440 39 39.5 Comparative Example 4 700 360 52 47.5 Comparative Example 5 645 305 62.5 55 Comparative Example 6 845 465 36 37 Comparative Example 7 835 435 46 47 Comparative Example 8 840 440 44 46 Comparative Example 9 835 430 45 47 Comparative Example 10 910 530 34 39

[0102] Table 5. Room temperature impact performance of Examples 1-10 and Comparative Examples 1-10

[0103]

[0104]

[0105] Table 6. High-temperature creep performance of Examples 1-10 and Comparative Examples 1-10 at 650°C

[0106]

[0107]

[0108] Table 7 Welding performance evaluation of Examples 1-10 and Comparative Examples 1-10

[0109] alloy Liquefaction cracks in the heat-affected zone Stress cracks in the heat-affected zone Example 1 none none Example 2 none none Example 3 none none Example 4 none none Example 5 none none Example 6 none none Example 7 none none Example 8 none none Example 9 none none Example 10 none none Comparative Example 1 none have Comparative Example 2 have none Comparative Example 3 none none Comparative Example 4 none none Comparative Example 5 none none Comparative Example 6 none have Comparative Example 7 none have Comparative Example 8 none have Comparative Example 9 none none Comparative Example 10 have have

[0110] As can be seen from the data in the table above, the nickel-based heat-resistant alloy prepared in the embodiments of the present invention has good mechanical properties. At room temperature: tensile strength R m ≥870MPa, yield strength R p0.2 ≥480MPa, reduction of area A≥40%, elongation after fracture ≥40%, impact energy A kv ≥110J; at a high temperature of 650℃: tensile strength R m ≥700MPa, yield strength R p0.2 The alloy exhibits a strength ≥380 MPa, a reduction of area (A) ≥41%, and an elongation at break ≥42%, along with excellent high-temperature creep rupture properties and weldability. Specifically, when the elements in the alloy satisfy the following relationships: 6.7 × Mo - 50.7 ≤ Co ≤ -1.65 × Mo + 25.3, 1.3 ≤ Al + Ti ≤ 1.7, 0.5 ≤ Ti / Al ≤ 0.8, and -15.6 < 3lgZr + 4lgB < -13.8, as in Examples 1–6, the alloy demonstrates even better overall performance.

[0111] Comparative Example 1 contained a higher amount of Mo, and Comparative Example 2 contained a higher amount of Co. Compared with Examples 1-10 of the present invention, the embodiments of the present invention use lower amounts of Co and Mo. The room temperature and high temperature instantaneous strength of the resulting alloys are basically equivalent to those of Comparative Examples 1 and 2, and the production cost of the alloys is reduced. Appropriately increasing the Co content is beneficial to improving the creep rupture performance of the alloy. However, when the Co content exceeds the solid solubility of Mo and Co, M6C carbides and intermetallic compound μ phases are easily formed, which reduces the creep rupture life of the alloy under low stress. Furthermore, excessive Mo and Co content leads to the formation of welding cracks, reducing the weldability of the alloy.

[0112] Compared with Comparative Example 3 and the Example, the addition of 0.2% element V slightly increased the strength of the alloy, but significantly reduced the plasticity of the alloy, and the impact energy at room temperature decreased to 105 J.

[0113] In Comparative Example 4, the Ti+Al content was only 1.15. Comparing Comparative Example 4, Examples 1-3, and Examples 7-9, it can be seen that the Ti+Al content significantly affects the material strength. As the Ti+Al content decreases, the material strength decreases. Ti+Al should be ≥1.2% to ensure the alloy's strength meets the requirements for use; preferably, Ti+Al ≥1.3%. The Ti / Al ratio has a certain impact on the alloy strength; the alloy strength increases with increasing Ti / Al ratio. When the Ti+Al content is similar, the creep rupture performance is similar. When the Ti / Al ratio is low, the absolute value of the mismatch increases, and γ′ coarsening is faster after long-term service, weakening the strengthening effect. Increasing the Ti / Al ratio is beneficial for improving creep rupture performance under low stress, but has less impact on creep rupture performance under high stress.

[0114] Comparative Examples 5 and 6 adjusted the amount of element C. As can be seen from Comparative Examples 5 and 6, the carbon content has a significant impact on the instantaneous strength at room temperature and high temperature. In Comparative Example 5, the carbon content was 0.041%, which is lower than 0.05%, resulting in insufficient strength of the alloy. The contribution of carbon content to strength has a certain threshold. In Comparative Example 6, the carbon content was 0.092%, which is higher than 0.08%. The increase in carbon content has a limited effect on improving strength, and excessively high carbon content causes a decrease in the plasticity of the alloy and the generation of welding cracks.

[0115] W was added to both Comparative Examples 7 and 8. In Comparative Example 7, the amount of W added was 0.41%, while in Comparative Example 8, in addition to the addition of 0.4% W, a high Co content of 11.87% was used. Adding a small amount of W can slightly improve the instantaneous strength of the alloy at room temperature and high temperature, but the improvement effect is not significant. Moreover, W promotes the formation of M6C, leading to the formation of welding cracks and reducing the weldability of the alloy.

[0116] Compared with Example 1, Comparative Example 9 did not add Zr. Since the addition of an appropriate amount of Zr can strengthen the grain boundaries, ZrC forms around Ti(C,N), which improves stress concentration, improves the strength and plasticity of the alloy, and increases the endurance strength of the alloy, the alloy prepared in Comparative Example 9 without the addition of Zr has reduced strength and plasticity, and its high-temperature endurance performance is also reduced to a certain extent.

[0117] In Example 1, Zr was added, and the SEM images of the carbonitrides in the alloy are shown below. Figure 3As shown, (a) is ZrC and (b) is clustered ZrC formed around Ti(C,N). ZrC forms around Ti(C,N), which is mostly rectangular. There is a large stress concentration at its tip and the matrix, which is easy to become a crack initiation site. In the alloy of the present invention, ZrC can precipitate in clusters around Ti(C,N), which effectively reduces stress concentration, reduces crack initiation sites, and improves the plasticity of the alloy.

[0118] In Comparative Example 10, excessive titanium was added, with Al+Ti ratio of 2.52 and Ti / Al ratio of 1.032. Although the resulting alloy exhibited high room temperature strength, high temperature strength at 650℃, and creep strength at 650℃, its plasticity and impact energy were significantly reduced. Furthermore, due to its high γ′ phase content, the forging deformation resistance was increased, making deformation more difficult. Moreover, the γ′ phase tends to coarsen during welding, leading to stress concentration and subsequent weld cracks, thus reducing the alloy's weldability.

[0119] (2) The room temperature mechanical properties and durability of Application Example 1 and Application Comparative Example 1 were tested, and the results are shown in Tables 8-9.

[0120] After the alloy was enlarged, segregation caused differences in mechanical properties at different locations. Samples were taken from the core, R / 2, and edge for testing. Specific sampling locations are shown below. Figure 4 As shown.

[0121] Table 8. Mechanical properties at room temperature and high temperature for Application Example 1 and Application Comparative Example 1

[0122]

[0123] Table 9. Durability at 650°C for Application Example 1 and Comparative Example 1

[0124]

[0125] In the same forging, segregation can lead to significant differences in mechanical properties at different locations, resulting in an alloy strength distribution where the edge > R / 2 > core. This invention employs a low-Co, low-Mo design and eliminates the addition of easily segregating elements such as W and V, thus reducing segregation in large-tonnage alloy forgings. Table 8 shows that in Application Example 1, the room temperature yield strength at R / 2 and the core is 96.6% and 95.9% of that at the edge, respectively, and the high-temperature yield strength at 650℃ is 97.7% and 96.6% of that at the edge, respectively. In Comparative Example 1, the room temperature yield strength at R / 2 and the core is 93.9% and 92.6% of that at the edge, respectively, and the high-temperature tensile yield strength at 650℃ is 93.2% and 90.9% of that at the edge, respectively. Furthermore, due to severe segregation in the core, Co, Mo, and W elements are more easily enriched in the core, forming brittle phases such as the μ phase, resulting in a decrease in core strength without an improvement in ductility and toughness. Table 9 shows that in Application Example 1, the strength at different locations 10 5 The extrapolated creep strength difference within hours was within 5 MPa, indicating good creep life; while in Comparative Example 1, the alloy 10 at different locations showed... 5 The difference in extrapolated endurance strength over hours is about 15 MPa, which is a relatively large deviation.

[0126] (3) SEM scans were performed on the alloys prepared in Example 1, Comparative Example 1, and Comparative Example 2. The results are as follows: Figure 5 , Figure 6 and Figure 7 As shown.

[0127] from Figure 5 The grain boundaries and intragranular M-type precipitates in the alloy prepared in Example 1 can be seen in the image. 23 C6 carbides; from Figure 6 The M6C carbide precipitated in the alloy prepared in Comparative Example 1 can be seen in the image; from Figure 7 The μ phase precipitated in the alloy prepared in Comparative Example 2 can be seen in the image. Figures 5-7 It can be seen that the microstructure of the alloy obtained by the present invention is based on M 23 C6 carbides are the main component. When the Mo content is high, M6C type carbides are easily formed, which can easily become crack initiation sites during endurance tests. When the Co content in the alloy is high, intermetallic compound μ phase is easily formed, which reduces the strength of the alloy and makes it prone to cracking.

[0128] In this invention, the terms "one embodiment," "some embodiments," "example," "specific example," or "some examples," etc., refer to a specific feature, structure, material, or characteristic described in connection with that embodiment or example, which is included in at least one embodiment or example of the invention. In this specification, the illustrative expressions of the above terms do not necessarily refer to the same embodiment or example. Furthermore, the specific features, structures, materials, or characteristics described may be combined in any suitable manner in one or more embodiments or examples. Moreover, without contradiction, those skilled in the art can combine and integrate the different embodiments or examples described in this specification, as well as the features of different embodiments or examples.

[0129] Although the above embodiments have been shown and described, it is understood that the above embodiments are exemplary and should not be construed as limiting the present invention. Any changes, modifications, substitutions and variations made to the above embodiments by those skilled in the art are within the protection scope of the present invention.

Claims

1. A nickel-based heat-resistant alloy, characterized in that, Includes: C: 0.05-0.08%, Cr: 21.0-23.0%, Co: 6.5-8.8%, Mo: 6.0-8.8%, Al: 0.70-1.40%, Ti: 0.40-0.80%, Nb: 0.15-0.55%, B: 0.002-0.005%, Zr: 0.015-0.06%, with the balance being nickel and unavoidable impurity elements, in mass percentage.

2. The nickel-based heat-resistant alloy according to claim 1, characterized in that, The impurity elements include at least one selected from Fe, Mn, Si, S, P, Cu, O, N, H, Pb, Sb, Sn, Bi, and As; the content of the impurities by mass percentage is: Fe≤0.5%, Mn≤0.10%, Si≤0.10%, S≤0.002%, P≤0.008%, Cu≤0.05%, N≤20ppm, O≤10ppm, H≤1ppm, Pb≤0.001%, Sb≤0.0025%, Sn≤0.0012%, Bi≤0.001%, and As≤0.002%.

3. The nickel-based heat-resistant alloy according to claim 1, characterized in that, The contents of Co and Mo satisfy the following relationship: 6.7×Mo-50.7≤Co≤-1.65×Mo+25.3, where Co and Mo refer to the mass percentages of elements Co and Mo in the nickel-based heat-resistant alloy after removing the percentage sign.

4. The nickel-based heat-resistant alloy according to claim 1, characterized in that, The contents of Al and Ti simultaneously satisfy the following relationships: 1.3≤Al+Ti≤1.7, 0.5≤Ti / Al≤0.8, where Al and Ti refer to the mass percentages of elements Al and Ti in the nickel-based heat-resistant alloy after removing the percentage sign.

5. The nickel-based heat-resistant alloy according to claim 1, characterized in that, The contents of B and Zr satisfy the following relationship: -15.6 < 3lgZr + 4lgB < -13.8, where B and Zr refer to the mass percentages of elements B and Zr in the nickel-based heat-resistant alloy after removing the percentage sign.

6. A method for preparing a nickel-based heat-resistant alloy according to any one of claims 1 to 5, characterized in that, Includes the following steps: (1) The raw materials are smelted, cast, homogenized and forged according to the design ratio to obtain a forging billet; (2) The forging billet obtained in step (1) is subjected to solution treatment and aging treatment to obtain a nickel-based heat-resistant alloy.

7. The method for preparing the nickel-based heat-resistant alloy according to claim 6, characterized in that, In step (1), the smelting is carried out using a process of vacuum induction melting + vacuum consumable remelting or vacuum induction melting + protective atmosphere electroslag melting + vacuum consumable remelting.

8. The method for preparing the nickel-based heat-resistant alloy according to claim 6 or 7, characterized in that, In step (1), the homogenization process is carried out using a graded heating process.

9. The method for preparing the nickel-based heat-resistant alloy according to claim 6 or 7, characterized in that, In step (2), the solution treatment includes heating the forging billet to 1150-1190°C and holding it for 60-90 minutes, followed by water cooling to room temperature; the aging treatment includes heating the forging to 700-750°C and holding it for 16-24 hours, followed by air cooling to room temperature.

10. The application of the nickel-based heat-resistant alloy according to any one of claims 1 to 5 or the nickel-based high-temperature alloy prepared by the preparation method according to any one of claims 6 to 9 in the rotor of an ultra-supercritical steam turbine.

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