High-temperature-resistant nickel-based wrought superalloy with high strength and toughness at 800 °c, preparation method thereof and turbine disc

Abstract

The application relates to the technical field of high-temperature alloys, in particular to a 800 DEG C-resistant high-strength and high-toughness nickel-based wrought high-temperature alloy, a preparation method thereof and a turbine disc. The nickel-based wrought high-temperature alloy comprises Cr 10-12%, Co 14-16%, W 2.5-3.5%, Mo 3.75-5%, Al 3.5-4.5%, Ti 2.5-3.1%, Nb 3-3.8%, C 0.01-0.08%, B 0.005-0.05%, V 0.01-1%, Zr 0.025-0.075%, Si 0.005-0.3%, Re 0.02-0.05%, Sc 0.001-0.08%, La 0.001-0.08%, Ce 0.001-0.05%, Nd 0.001-0.05%, Mg 0.001-0.15%, and the balance is Ni. The application realizes good hot working performance of the alloy in engineering by reasonably matching the components of the nickel-based high-temperature alloy and controlling the coarsening of the gamma prime phase through a process, and the turbine disc meeting the requirements can be obtained through heat treatment.

Description

Technical Field

[0001] This invention relates to the field of high-temperature alloy technology, and in particular to a high-strength and high-toughness nickel-based deformable high-temperature alloy resistant to 800℃, its preparation method, and a turbine disk. Background Technology

[0002] Turbine disks are crucial hot-end components of aero-engines, operating in harsh environments and primarily constructed from wrought superalloys. With the development of the aviation industry, the exit temperature of aero-engines is gradually increasing, thus requiring the nickel-based wrought superalloys used in turbine disks to possess superior mid-temperature (600–800℃) service performance. The strengthening and toughening design methods for wrought superalloys mainly include: precipitation strengthening through the addition of Al, Ti, and Nb to form γ” or γ' phases; solid solution strengthening through the addition of W and Mo; and grain boundary strengthening through the addition of microalloying elements such as B, Zr, and C.

[0003] Generally, there are two main process routes for high-alloy superalloys: powder metallurgy, which reduces the severity of segregation, but it is difficult to solve defects such as original particle boundaries, inclusions and thermally induced voids, which seriously affect the mechanical properties of powder superalloys, such as low-cycle fatigue; and casting-forging process, which is the current mainstream process, mature and suitable for industrial application.

[0004] Currently, the main typical grades of nickel-based wrought superalloys include GH4169 alloy (resistant to 650℃), GH169D alloy (resistant to 700℃), GH4720Li alloy (resistant to 750℃), and GH4065A alloy. The degree of alloying in these alloys is gradually increasing. The high degree of alloying will significantly reduce the plasticity of the alloy in processing (high temperature) and service (medium temperature) environments. The increasingly high degree of alloying leads to severe element segregation in high-temperature alloy ingots, complex solidification segregated phases, severe element segregation in large-size alloy ingots, difficulty in eliminating segregated structure inheritance, ingot size limitation to small ingots, narrow hot working window, high deformation resistance, and easy cracking. These problems make the billeting process and microstructure control extremely difficult, resulting in extremely low yield and uneven microstructure. These issues have become key problems restricting the preparation and application of high-alloyed high-temperature alloy high-quality bars.

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

[0006] One objective of this invention is to provide a high-strength and high-toughness nickel-based wrought high-temperature alloy that can withstand temperatures up to 800°C. This alloy not only possesses both strength and toughness in the mid-temperature range but also exhibits high-temperature machinability, thus providing the necessary high-thermal-strength nickel-based wrought high-temperature material for high thrust-ratio engines.

[0007] Another object of the present invention is to provide a method for preparing a high-strength and high-toughness nickel-based deformable high-temperature alloy resistant to 800℃.

[0008] Another object of the present invention is to provide a turbine disk made of the above-mentioned high-strength and high-toughness nickel-based high-temperature alloy that can withstand 800℃ deformation.

[0009] To achieve the above-mentioned objectives of the present invention, one aspect of the present invention provides a high-strength and high-toughness nickel-based high-temperature alloy resistant to deformation at 800℃, comprising the following components by mass percentage:

[0010] Cr 10%–12%, Co 14%–16%, W 2.5%–3.5%, Mo 3.75%–5%, Al 3.5%–4.5%, Ti 2.5%–3.1%, Nb 3%–3.8%, C 0.01%–0.08%, B 0.005%–0.05%, V 0.01%–1%, Zr 0.025%–0.075%, Si 0.005%–0.3%, Re 0.02%–0.05%, Sc 0.001%–0.08%, La 0.001%–0.08%, Ce 0.001%–0.05%, Nd 0.001%–0.05%, Mg 0.001%–0.15%, with the balance being Ni and unavoidable impurities.

[0011] In a specific embodiment of the present invention, the sum of the mass percentages of Al, Ti and Nb satisfies: Al+Ti+Nb≥9.5%; the mass ratio of Al to Ti, Al / Ti, is 1.15 to 1.45.

[0012] In a specific embodiment of the present invention, the sum of the mass percentages of Mo and W satisfies: Mo + W ≥ 7.5%; the mass ratio of Mo to W, Mo / W, is 1.3 to 1.45.

[0013] In a specific embodiment of the present invention, the sum of the mass percentages of C, B, and Zr satisfies: C + B + Zr ≤ 0.15%.

[0014] In a specific embodiment of the present invention, the sum of the mass percentages of Sc, La, Ce, and Nd satisfies: Sc + La + Ce + Nd ≤ 0.2%.

[0015] In a specific embodiment of the present invention, the unavoidable impurities include any one or more of S, P, N, O, Cu, Ag, Si, Mn, Ca, Sn, and Pb. Further, in the alloy, by mass percentage, the impurity elements include: S ≤ 0.0015%, P ≤ 0.015%, N ≤ 0.003%, O ≤ 0.0015%, Cu ≤ 0.10%, Ag ≤ 0.0005%, Si ≤ 0.10%, Mn ≤ 0.4%, Ca ≤ 0.005%, Sn ≤ 0.005%, and Pb ≤ 0.001%.

[0016] Another aspect of the present invention provides a method for preparing any one of the above-described high-strength and high-toughness nickel-based deformable superalloys resistant to 800℃, comprising the following steps:

[0017] (a) According to the alloy composition, alloy ingots are obtained through a three-stage smelting process;

[0018] (b) The alloy ingot is subjected to multi-stage homogenization annealing at 1100-1220°C to obtain an annealed ingot;

[0019] (c) The annealed ingot is upset at 1060-1180°C and then radially forged at 1020-1150°C.

[0020] In a specific embodiment of the present invention, the multi-stage homogenization annealing includes: holding the alloy ingot at 490-510°C for more than 4 hours, then raising the temperature to 990-1010°C and holding it for more than 8 hours, then raising the temperature to 1110-1140°C and holding it for more than 15 hours, then raising the temperature to 1150-1180°C and holding it for more than 25 hours, then raising the temperature to 1190-1220°C and holding it for more than 30 hours, and then furnace cooling.

[0021] In a specific embodiment of the present invention, a high-speed forging mill is used for the upsetting and drawing of the billet, with an upsetting deformation of 10% to 55% per forging and a drawing deformation of 15% to 60% per forging. Furthermore, in the upsetting and drawing of the billet, the forging temperature is gradually reduced in the range of 1060 to 1180°C per forging.

[0022] In a specific embodiment of the present invention, in the upsetting and drawing of the billet, the time for heat preservation in the furnace after each deformation is 1 to 6 hours.

[0023] In a specific embodiment of the present invention, the radial forging process employs a multi-pass deformation method to prepare an alloy of predetermined specifications. Further, in the radial forging process, the holding time at 1020–1150°C is 30–300 min.

[0024] In a specific embodiment of the present invention, the method further includes: heat treatment of the alloy after radial forging, wherein the heat treatment includes solution treatment and aging treatment. Further, the solution treatment includes: holding at 1100–1150°C for 2–6 hours; the aging treatment includes: holding at 800–900°C for 2–6 hours followed by air cooling, and then holding at 750–820°C for 8–24 hours followed by air cooling.

[0025] In a specific embodiment of the present invention, the average grain size of the heat-treated alloy is grade 8 or higher, with a grade difference of <2. Further, the γ' phase content of the heat-treated alloy is 50wt% to 58wt%.

[0026] In a specific embodiment of the present invention, the tensile properties of the heat-treated alloy satisfy the following:

[0027] Tensile strength at room temperature ≥1510MPa, yield strength at room temperature ≥1130MPa;

[0028] Tensile strength at 750℃ ≥1110MPa, yield strength at 750℃ ≥960MPa;

[0029] Tensile strength at 800℃ ≥1000MPa, yield strength at 800℃ ≥900MPa.

[0030] In a specific embodiment of the present invention, the creep rupture life of the heat-treated alloy satisfies:

[0031] The service life at 750℃ / 620MPa is ≥60h;

[0032] Duration life at 800℃ / 500MPa ≥45h.

[0033] In another aspect, the present invention provides a turbine disk made of any of the above-mentioned high-strength and high-toughness nickel-based high-temperature alloys that can withstand 800℃ deformation.

[0034] Compared with the prior art, the beneficial effects of the present invention are as follows:

[0035] (1) This invention designs the composition of nickel-based high-temperature alloys in a reasonable way, adds certain rare earth elements and trace elements such as C, B, and Zr, improves the grain boundary bonding force of the alloy, reduces or eliminates the tendency of brittleness in the service temperature range, and coarsens the γ' phase through process control, so that the high-temperature alloy can be processed and prepared, avoiding the mid-temperature brittleness and low plasticity of the alloy in the service temperature range, so that the alloy can take into account both high strength and toughness and good process performance.

[0036] (2) In combination with the alloy composition, the present invention adopts a multi-stage homogenization annealing process, upsetting and radial forging combined billet opening process to prepare a high homogeneous and low segregation alloy and ensure its high strength and toughness at 800℃; through process control, the strengthening effect of the γ' phase is weakened, and the alloy is controlled into a fine grain structure, so that the alloy can be hot-processed in engineering.

[0037] (3) The nickel-based wrought high-temperature alloy of the present invention meets the high thermal strength requirements of high thrust ratio engines, which is of great significance to the development of the aviation field. Attached Figure Description

[0038] To more clearly illustrate the specific embodiments of the present invention or the technical solutions in the prior art, the drawings used in the description of the specific embodiments or the prior art will be briefly introduced below. Obviously, the drawings described below are some embodiments of the present invention. For those skilled in the art, other drawings can be obtained from these drawings without creative effort.

[0039] Figure 1 The alloy system of this invention exhibits non-equilibrium solidification segregation characteristics such as alloy segregation phases and elemental segregation.

[0040] Figure 2 The multi-stage high-temperature diffusion homogenization annealing temperature curve provided in the embodiments of the present invention;

[0041] Figure 3 This is a grain structure diagram of the rod material obtained in Example 1 of the present invention;

[0042] Figure 4 The grain structure diagram is shown for the rod material prepared in Comparative Example 1.

[0043] Figure 5 The image shows the microstructure of the annealed alloy ingot obtained in Comparative Example 1.

[0044] Figure 6 This is a tensile fracture morphology image of the bar obtained in Example 1 of the present invention at 800℃.

[0045] Figure 7 The image shows the fracture morphology of a long-term specimen of the bar material prepared in Example 1 of this invention under an 800℃ / 500MPa environment.

[0046] Figure 8 The image shows the fracture surface morphology of the bar obtained in Comparative Example 1 under 800℃ tensile testing.

[0047] Figure 9 The image shows the fracture morphology of the bar obtained in Comparative Example 1 under 800℃ / 500MPa conditions. Detailed Implementation

[0048] The technical solution of the present invention will be clearly and completely described below with reference to the accompanying drawings and specific embodiments. However, those skilled in the art will understand that the embodiments described below are some embodiments of the present invention, but not all embodiments, and are only used to illustrate the present invention, and should not be regarded as limiting the scope of the present invention. Based on the embodiments of the present invention, all other embodiments obtained by those skilled in the art without creative effort are within the scope of protection of the present invention. Where specific conditions are not specified in the embodiments, conventional conditions or conditions recommended by the manufacturer shall be followed. Where the manufacturers of reagents or instruments are not specified, they are all conventional products that can be purchased commercially.

[0049] This invention provides a high-strength, high-toughness nickel-based high-temperature alloy resistant to deformation at 800℃, comprising the following components by mass percentage:

[0050] Cr 10%–12%, Co 14%–16%, W 2.5%–3.5%, Mo 3.75%–5%, Al 3.5%–4.5%, Ti 2.5%–3.1%, Nb 3%–3.8%, C 0.01%–0.08%, B 0.005%–0.05%, V 0.01%–1%, Zr 0.025%–0.075%, Si 0.005%–0.3%, Re 0.02%–0.05%, Sc 0.001%–0.08%, La 0.001%–0.08%, Ce 0.001%–0.05%, Nd 0.001%–0.05%, Mg 0.001%–0.15%, with the balance being Ni and unavoidable impurities.

[0051] This invention improves the hot working and mechanical properties of nickel-based superalloys through a rational design of their composition. The alloy of this invention is a highly alloyed, difficult-to-deform nickel-based superalloy. The γ' phase content in the alloy is approximately 10% higher than that in alloys with a 750℃ resistance rating. Simultaneously, the content of solid solution strengthening elements is also relatively high, achieving excellent high-temperature strength at 800℃. The alloy incorporates rare earth elements such as Sc and grain boundary strengthening elements such as Zr, improving grain boundary bonding strength and shifting the isohyetal transition temperature between grain boundaries and the grain interior to higher temperatures, resulting in better toughness and preventing mid-temperature brittleness around 800℃. Combined with process control during hot working, the strengthening effect of the γ' phase is weakened, resulting in a fine-grained structure that allows for hot working in engineering applications.

[0052] The alloy of this invention contains a large number of solid solution strengthening elements, precipitation strengthening elements, grain boundary strengthening elements and rare earth elements, and its composition characteristics are as follows.

[0053] The addition of precipitation-strengthening elements Al, Ti, and Nb to the alloy results in a γ' phase content of 50wt%–58wt%, significantly higher than other high-heat-strength nickel-based wrought superalloys. This abundant precipitation-strengthening phase ensures sufficient strength. Nb enters the γ' phase to form Ni3(Al, Ti, Nb), increasing the volume fraction of the γ' phase and the reverse domain boundary energy, thereby improving the alloy's strength. To balance strength and workability, this invention further optimizes the sum of the mass percentages of Al, Ti, and Nb to satisfy: Al + Ti + Nb ≥ 9.5%, and the mass ratio of Al to Ti (Al / Ti) is 1.15–1.45.

[0054] In different embodiments, the mass percentage of Al in this invention can be 3.5%, 3.6%, 3.7%, 3.8%, 3.9%, 4%, 4.1%, 4.2%, 4.3%, 4.4%, 4.5%, or any combination thereof; the mass percentage of Ti can be 2.5%, 2.6%, 2.8%, 2.9%, 3%, 3.1%, or any combination thereof; and the mass percentage of Nb can be 3%, 3.1%, 3.2%, 3.3%, 3.4%, 3.5%, 3.6%, 3.7%, 3.8%. Furthermore, the sum of the mass percentages of Al, Ti, and Nb is 9.5% to 10.5%, such as 9.5%, 9.6%, 9.7%, 9.8%, 9.9%, 10%, 10.1%, 10.2%, 10.3%, 10.4%, 10.5%, or any combination thereof; the mass ratio of Al to Ti, Al / Ti, can be 1.15, 1.2, 1.25, 1.3, 1.35, 1.4, 1.45, or any combination thereof.

[0055] Adding solid solution strengthening elements W and Mo to alloys causes significant lattice distortion and reduces stacking fault energy, thereby improving the alloy's yield strength and creep properties. However, adding excessive or improperly proportioned W and Mo can lead to the formation of a harmful brittle TCP phase during long-term use. Under thermo-mechanical coupling, cracks initiate and propagate at the TCP phase-matrix interface, resulting in deterioration of the alloy's properties. Therefore, this invention further controls the sum of the mass percentages of Mo and W to satisfy: W + Mo ≥ 7.5%, and the mass ratio of M to W, Mo / W, is 1.3–1.45, to ensure sufficient solid solution strengthening effect in the alloy.

[0056] In different embodiments, the mass percentage of W in this invention can be 2.5%, 2.6%, 2.7%, 2.8%, 2.9%, 3%, 3.1%, 3.2%, 3.3%, 3.4%, 3.5%, or any two of these components; the mass percentage of Mo can be 3.75%, 3.9%, 4%, 4.1%, 4.2%, 4.3%, 4.4%, 4.5%, 4.6%, 4.7%, 4.8%, 4.9%, or 5%. Furthermore, the sum of the mass percentages of Mo and W is 7.5% to 8.2%, such as a range of 7.5%, 7.6%, 7.7%, 7.8%, 7.9%, 8%, 8.1%, 8.2%, or any combination thereof; the mass ratio of Mo to W, Mo / W, can be a range of 1.3, 1.32, 1.35, 1.38, 1.4, 1.42, 1.45, or any combination thereof.

[0057] The addition of carbon (C) to the alloy of this invention causes the element to segregate to grain boundaries or form grain boundary carbides, which precipitate at the grain boundaries in a blocky or granular form. This inhibits abnormal grain growth, prevents grain boundary sliding, strengthens grain boundaries, and improves creep strength and ductility. However, excessive addition can lead to continuous precipitation of carbides along the grain boundaries in a thin film, causing grain boundary embrittlement and notch sensitivity in the alloy. Therefore, the C content in this invention is controlled within the range of 0.01% to 0.08%. In different embodiments, the mass percentage of C can be 0.01%, 0.02%, 0.03%, 0.04%, 0.05%, 0.06%, 0.07%, 0.08%, or any combination thereof.

[0058] The alloy of this invention incorporates Zr and B elements, which segregate at grain boundaries, enhancing grain boundary bonding, reducing interfacial energy, improving grain boundary morphology, and increasing creep rupture strength and low-cycle fatigue performance. Simultaneously, Zr acts as a purifying agent, reacting with carbides or carbides of elements such as C and S in the alloy to remove harmful impurities like S and P. However, excessive addition of B can lead to the precipitation of borides at grain boundaries or induce thermal cracking. Therefore, in this invention, the Zr content is controlled at 0.025%–0.075%, and the B content at 0.005%–0.05%. In different embodiments, the mass percentage of Zr can be 0.025%, 0.035%, 0.045%, 0.055%, 0.065%, 0.075%, or any combination thereof, and the mass percentage of B can be 0.005%, 0.008%, 0.01%, 0.02%, 0.03%, 0.04%, 0.05%, or any combination thereof.

[0059] The present invention further optimizes the sum of the mass percentages of C, B, and Zr to satisfy: C + B + Zr ≤ 0.15%, so as to improve grain boundary bonding, increase creep strength and ductility, while reducing or eliminating the tendency for brittleness in the service temperature range, and avoiding grain boundary embrittlement and hot cracking. In different embodiments, the sum of the mass percentages of C, B, and Zr can be 0.045%, 0.06%, 0.07%, 0.08%, 0.09%, 0.1%, 0.11%, 0.12%, 0.13%, 0.14%, 0.15%, or any combination thereof.

[0060] The addition of rare earth elements Sc, La, Ce, and Nd to the alloy of this invention can reduce the critical undercooling for dendrite nucleation, generating compositional undercooling at the solid-liquid interface front during alloy solidification, reducing the critical nucleation radius and nucleation work, increasing the nucleation rate, thereby refining dendrites and effectively reducing segregation. Simultaneously, rare earth elements can remove oxygen and purify grain boundaries, reducing or preventing intergranular cracking in the mid-temperature range. Rare earth elements can also improve the alloy's oxidation resistance. However, excessive rare earth elements are prone to oxidation, and the rare earth-rich phase is difficult to eliminate during subsequent homogenization and hot deformation processes. Therefore, the rare earth content in this invention should be strictly controlled. Specifically, the Sc content is controlled at 0.001%–0.08%, the La content at 0.001%–0.08%, the Ce content at 0.001%–0.05%, and the Nd content at 0.001%–0.05%. In different embodiments, the mass percentage of Sc can be a range of 0.001%, 0.005%, 0.01%, 0.02%, 0.04%, 0.06%, 0.08%, or any combination thereof; the mass percentage of La can be a range of 0.001%, 0.005%, 0.01%, 0.02%, 0.04%, 0.06%, 0.08%, or any combination thereof; the mass percentage of Ce can be a range of 0.001%, 0.005%, 0.01%, 0.02%, 0.03%, 0.04%, 0.05%, or any combination thereof; and the mass percentage of Nd can be a range of 0.001%, 0.005%, 0.01%, 0.02%, 0.03%, 0.04%, 0.05%, or any combination thereof.

[0061] Further adjustments can be made to ensure that the sum of the mass percentages of Sc, La, Ce, and Nd satisfies: Sc + La + Ce + Nd ≤ 0.2%, such as 0.1% to 0.2%, to optimize the performance of each component. In different embodiments, the sum of the mass percentages of Sc, La, Ce, and Nd can be within the range of 0.1%, 0.11%, 0.12%, 0.13%, 0.14%, 0.15%, 0.16%, 0.17%, 0.18%, 0.19%, 0.2%, or any combination thereof.

[0062] Cr, the main carbide-forming element in the alloy of this invention, also enhances the alloy's oxidation and corrosion resistance. Too low a Cr content leads to a decrease in the alloy's oxidation and corrosion resistance, and also results in the formation of the α-phase, which negatively impacts the alloy's ductility and toughness. Therefore, this invention controls the Cr content in the alloy to be between 10% and 12%, combined with the remaining elements, to achieve a balance between improved oxidation and corrosion resistance as well as enhanced ductility and toughness. In different embodiments, the mass percentage of Cr in this invention can be within the range of 10%, 10.2%, 10.5%, 10.8%, 11%, 11.2%, 11.5%, 11.8%, 12%, or any combination thereof.

[0063] In specific embodiments of the present invention, unavoidable impurities include any one or more of S, P, N, O, Cu, Ag, Si, Mn, Ca, Sn, and Pb. Further, in the alloy, by mass percentage, the impurity elements include: S ≤ 0.0015%, P ≤ 0.015%, N ≤ 0.003%, O ≤ 0.0015%, Cu ≤ 0.10%, Ag ≤ 0.0005%, Si ≤ 0.10%, Mn ≤ 0.4%, Ca ≤ 0.005%, Sn ≤ 0.005%, and Pb ≤ 0.001%.

[0064] In a specific embodiment of the present invention, the high-strength and high-toughness nickel-based high-temperature alloy resistant to 800℃ deformation comprises the following components by mass percentage:

[0065] Cr 10.5%~11.5%, Co 15%~15.5%, W 3.2%~3.5%, Mo 4.4%~4.6%, Al3.7%~3.9%, Ti 2.7%~2.8%, Nb 3.3%~3.5%, C 0.04%~0.06%, B0.005%~0.006%, V 0.5%~0.7%, Zr 0.05%~0.06%, Si 0.015%~0.03%, Re 0.02%~0.03%, Sc 0.03%~0.06%, La 0.02%~0.03%, Ce 0.02%~0.03%, Nd 0.03%~0.05%, Mg 0.03% to 0.04%, with the balance being Ni and unavoidable impurities.

[0066] Another aspect of the present invention provides a method for preparing any of the above-mentioned high-strength and high-toughness nickel-based deformable high-temperature alloys resistant to 800℃, comprising the following steps:

[0067] (a) According to the alloy composition, alloy ingots are obtained through a three-stage smelting process;

[0068] (b) The alloy ingot is subjected to multi-stage homogenization annealing at 1100-1220℃ to obtain the annealed ingot;

[0069] (c) The annealed ingot is upset at 1060-1180°C and then radially forged at 1020-1150°C.

[0070] The material of this invention has a high degree of alloying and requires high purity of the alloy. By using a triple smelting process of vacuum induction melting + electroslag remelting + vacuum induction remelting, combined with the alloy composition of this invention, ingots with high purity can be obtained.

[0071] In a specific embodiment of the present invention, during the vacuum induction melting of the triple smelting process, rare earth elements Sc, La, Ce, and Nd are added in the form of a master alloy; wherein the master alloy is an Al-rare earth master alloy or a Ni-rare earth master alloy. For example, they can be added in the form of Al-Sc, Al-La, Al-Ce, and Al-Nd, or in the form of Ni-Sc, Ni-La, Ni-Ce, and Ni-Nd, to improve the yield of rare earth elements.

[0072] In a specific embodiment of the present invention, the intermediate alloy is added during the low-temperature refining stage (1450-1580°C) to improve the yield of rare earth elements.

[0073] In a specific embodiment of the present invention, the triple smelting process includes:

[0074] A: According to the alloy composition, the alloy is smelted in a vacuum induction furnace, the full melting temperature is controlled at 1510~1580℃, the refining temperature is controlled at 1510~1580℃, and the vacuum induction electrode rod is obtained by casting at 1420~1480℃.

[0075] B: The vacuum induction electrode rod obtained in step A is subjected to electroslag remelting under a protective atmosphere to obtain an electroslag remelted ingot; argon gas is used for protection throughout step B, and the argon gas flow rate is 100-200 L / min.

[0076] C: The electroslag remelted ingot obtained in step B is subjected to vacuum arc remelting to obtain a vacuum arc remelted ingot; the vacuum arc remelting rate is 2.7 to 3.8 kg / min.

[0077] In a specific embodiment of the present invention, the multi-stage homogenization annealing includes: holding the alloy ingot at 490-510°C for more than 4 hours, then raising the temperature to 990-1010°C and holding it for more than 8 hours, then raising the temperature to 1110-1140°C and holding it for more than 15 hours, then raising the temperature to 1150-1180°C and holding it for more than 25 hours, then raising the temperature to 1190-1220°C and holding it for more than 30 hours, and then furnace cooling.

[0078] Figure 1 To illustrate the non-equilibrium solidification segregation characteristics of the alloy system in this invention (such as Example 1), including alloy segregation phases and elemental segregation, this invention proposes a specific multi-stage homogenization annealing process based on these characteristics. (Refer to...) Figure 2 The temperature curves are designed to eliminate non-equilibrium segregation and / or elemental segregation at each temperature range, thereby obtaining a homogenized, low-segregation alloy ingot. Annealing at 1110–1140℃ and 1150–1180℃ helps eliminate low-melting-point phases, while annealing at 1190–1220℃ helps eliminate ingot compositional segregation and improves the hot working properties of the alloy.

[0079] In different embodiments, multi-stage homogenization annealing may include: holding the alloy ingot at 490℃, 495℃, 500℃, 505℃, or 510℃ for more than 4 hours, then raising the temperature to 990℃, 995℃, 1000℃, 1005℃, or 1010℃ and holding for more than 8 hours, and then raising the temperature to 1110℃, 1115℃, 1120℃, 1125℃, 1130℃, or 1115℃. Heat to 135℃ or 1140℃ and hold for more than 15 hours, then raise the temperature to 1150℃, 1155℃, 1160℃, 1165℃, 1170℃, 1175℃ or 1180℃ and hold for more than 25 hours, then raise the temperature to 1190℃, 1195℃, 1200℃, 1205℃, 1210℃, 1215℃ or 1220℃ and hold for more than 30 hours, then cool in the furnace.

[0080] In actual operation, the heating rate between adjacent insulation temperatures can be 2 to 120℃ / h. For example, it can be a range of 2℃ / h, 10℃ / h, 20℃ / h, 40℃ / h, 50℃ / h, 60℃ / h, 80℃ / h, 100℃ / h, 120℃ / h, or any combination thereof.

[0081] In a specific embodiment of the present invention, the heating rate from room temperature to 490-510°C can be 85-95°C / h; the heating rate from 490-510°C to 990-1010°C can be 75-85°C / h; the heating rate from 990-1010°C to 1110-1140°C can be 45-55°C / h; the heating rate from 1110-1140°C to 1150-1180°C can be 8-12°C / h; and the heating rate from 1150-1180°C to 1190-1220°C can be 3-7°C / h.

[0082] In a specific embodiment of the present invention, a high-speed forging mill is used for upsetting and drawing the billet, with an upsetting deformation of 10% to 55% per heat and a drawing deformation of 15% to 60% per heat. Furthermore, during upsetting and drawing, the forging temperature is gradually reduced in each heat within the range of 1060 to 1180°C.

[0083] In practice, when using a high-speed forging machine for upsetting and drawing, the annealed ingot is heated to 1060-1180℃, wrapped with insulation cotton, and after the insulation is completed, it is taken out of the furnace for forging. The high-speed forging machine is used to repeatedly upset and draw the ingot.

[0084] In different embodiments, the deformation per upsetting pass can be 10%, 20%, 30%, 40%, 50%, 55%, or any combination thereof; the elongation deformation per drawing pass can be 15%, 20%, 30%, 40%, 50%, 60%, or any combination thereof.

[0085] In a specific embodiment of the present invention, during the upsetting and drawing process, the time for reheating in the furnace after each deformation is 1 to 6 hours. Reheating in the furnace after each deformation can achieve static recrystallization, refine grains, and improve thermoplasticity.

[0086] In a specific embodiment of the present invention, during the upsetting and drawing of the billet, the forging temperature can be gradually reduced in three forging temperature ranges; wherein, the first forging temperature is 1170-1180℃, the second forging temperature is 1150-1160℃, and the third forging temperature is 1130-1140℃.

[0087] In practice, the final forging temperature of upset drawing billet is ≥950℃.

[0088] In the forging process of the present invention, which involves progressively reducing the forging temperature, the forging temperature range is a two-phase region. During the forging process, the γ' phase gradually coarsens to the micrometer level. The micrometer-level γ' phase reduces the deformation resistance of the alloy, promotes the dynamic recrystallization of the alloy, and thus prepares fine-grained rods, improving the thermoplasticity of the alloy.

[0089] In a specific embodiment of the present invention, during the upsetting and drawing process, the upsetting deformation gradually increases, that is, the upsetting deformation at the first forging temperature is less than the upsetting deformation at the second forging temperature, and the upsetting deformation at the second forging temperature is less than the upsetting deformation at the third forging temperature.

[0090] In a specific embodiment of the present invention, the radial forging process employs a multi-pass deformation method to prepare an alloy of predetermined specifications. Furthermore, in the radial forging process, the holding time at 1020–1150°C is 30–300 min.

[0091] In practice, before radial forging the intermediate billet after upsetting, the hammer of the radial forging machine is preheated at 350–650℃ for ≥4 hours. During radial forging, the intermediate billet is returned to the furnace and encased, then held at 1020℃, 1040℃, 1050℃, 1060℃, 1080℃, 1100℃, 1120℃, or 1150℃ for 30 min, 60 min, 90 min, 120 min, 150 min, 180 min, 240 min, or 300 min respectively on the radial forging machine. The billet is then forged into bars of the predetermined specifications in a single pass on the radial forging machine and air-cooled to room temperature. This radial forging process ensures sufficient deformation of the outer edge of the bar, achieving dynamic recrystallization to obtain an equiaxed grain structure.

[0092] In a specific embodiment of the present invention, the process further includes: heat treatment of the alloy after radial forging, the heat treatment including solution treatment and aging treatment. Further, the solution treatment includes: holding at 1100–1150°C for 2–6 hours; the aging treatment includes: holding at 800–900°C for 2–6 hours followed by air cooling, and then holding at 750–820°C for 8–24 hours followed by air cooling.

[0093] During the solution treatment, the γ' phase is fully re-dissolved. Combined with the subsequent first-stage aging treatment, a γ' phase with a size of 100–500 μm can be formed. After the second-stage aging treatment, a uniformly dispersed γ' phase with a size of 10–50 μm can be formed.

[0094] In a specific embodiment of the present invention, the average grain size of the heat-treated alloy is grade 8 or higher, with a grade difference of <2. Further, the γ' phase content of the heat-treated alloy is 50wt% to 58wt%.

[0095] The nickel-based wrought superalloys prepared by the method of this invention can achieve an average grain size of grade 8 or higher, such as grade 8, grade 8.5, grade 9, grade 9.5, or any combination thereof. The grain size difference at various locations in the bar is <2 grades, for example, the grain size difference at the center, R / 2, and edge is <2 grades, such as grade 1.5, grade 1, grade 0.5, grade 0, or any combination thereof.

[0096] The alloy of the present invention ensures a high γ' phase content, thereby guaranteeing sufficient strength. In different embodiments, the γ' phase content can be a range of 50 wt%, 51 wt%, 52 wt%, 53 wt%, 54 wt%, 55 wt%, 56 wt%, 57 wt%, 58 wt%, or any combination thereof.

[0097] In a specific embodiment of the present invention, the tensile properties of the heat-treated alloy satisfy the following:

[0098] Tensile strength at room temperature ≥1510MPa, yield strength at room temperature ≥1130MPa;

[0099] Tensile strength at 750℃ ≥1110MPa, yield strength at 750℃ ≥960MPa;

[0100] Tensile strength at 800℃ ≥1000MPa, yield strength at 800℃ ≥900MPa.

[0101] In different embodiments, the tensile properties of the heat-treated alloy can be exemplarily described as follows:

[0102] The room temperature tensile strength can be in the range of 1510MPa, 1530MPa, 1550MPa, 1580MPa, 1600MPa, 1620MPa, 1650MPa, 1680MPa or any combination thereof; the room temperature yield strength can be in the range of 1130MPa, 1150MPa, 1180MPa, 1200MPa, 1220MPa, 1255MPa or any combination thereof.

[0103] The tensile strength at 750℃ can be 1110MPa, 1150MPa, 1160MPa, 1180MPa, 1200MPa, 1220MPa, 1240MPa, 1260MPa, 1270MPa, 1275MPa, or any combination thereof; the yield strength at 750℃ can be 960MPa, 980MPa, 1000MPa, 1020MPa, 1050MPa, 1080MPa, 1090MPa, or any combination thereof.

[0104] The tensile strength at 800℃ can be in the range of 1000MPa, 1020MPa, 1050MPa, 1080MPa, 1100MPa, 1120MPa, 1130MPa or any combination thereof; the yield strength at 800℃ can be in the range of 900MPa, 910MPa, 920MPa, 930MPa, 940MPa, 950MPa or any combination thereof.

[0105] In a specific embodiment of the present invention, the creep rupture life of the heat-treated alloy satisfies:

[0106] The creep life at 750℃ / 620MPa is ≥60h;

[0107] Duration life at 800℃ / 500MPa ≥45h.

[0108] In different embodiments, the creep rupture life of the heat-treated alloy can be exemplarily exemplified as follows:

[0109] The creep rupture life at 750℃ / 620MPa can be 60h, 70h, 80h, 100h, 110h, 120h, 130h, 140h, 150h, 160h, 164h or any combination thereof.

[0110] The creep life at 800℃ / 500MPa can be a range of 45h, 50h, 60h, 70h, 80h, 85h, or any combination thereof.

[0111] In another aspect, the present invention provides a turbine disk made of any of the aforementioned nickel-based wrought high-temperature alloys.

[0112] Example 1

[0113] This embodiment provides a nickel-based wrought superalloy comprising the following components by mass percentage: Cr 10.5%, Co 15.0%, W 3.4%, Mo 4.5%, Al 3.9%, Ti 2.8%, Nb 3.3%, C 0.06%, B 0.005%, V 0.6%, Zr 0.05%, Si 0.015%, Re 0.03%, Sc 0.06%, La 0.02%, Ce 0.02%, Nd 0.05%, Mg 0.03%, P 0.015%, S 0.0006%, and Ni balance.

[0114] The preparation method of nickel-based wrought superalloys includes the following steps:

[0115] (1) Prepare a consumable ingot with a specification of Φ508mm by a triple smelting process of vacuum induction melting + protective atmosphere electroslag remelting + vacuum consumable remelting according to the alloy composition. Polish the surface of the consumable ingot and spray it with conventional anti-oxidation coating.

[0116] During vacuum induction melting, rare earth elements Sc, La, Ce, and Nd are added in the form of intermediate alloys Ni-Sc, Ni-La, Ni-Ce, and Ni-Nd respectively during the low-temperature refining stage at 1480–1510℃; the full melting temperature is controlled at 1520–1550℃, the refining temperature at 1520–1540℃, and the casting is carried out at 1440–1450℃.

[0117] During electroslag remelting under a protective atmosphere, the argon flow rate is 150±10 L / min;

[0118] During vacuum arc remelting, the vacuum arc remelting rate is 3.2 kg / min.

[0119] (2) The self-consumable ingot obtained in step (1) is subjected to multi-stage homogenization annealing to obtain an annealed ingot with a length of L. The multi-stage homogenization annealing includes: heating to 500℃ at 90℃ / h and holding for 4h, then heating to 1000℃ at 80℃ / h and holding for 8h, then heating to 1140℃ at 50℃ / h and holding for 15h, then heating to 1180℃ at 10℃ / h and holding for 25h, then heating to 1210℃ at 5℃ / h and holding for 30h, and then furnace cooling.

[0120] (3) The annealed ingot obtained in step (2) is subjected to upsetting and forging, specifically including the following steps:

[0121] ① The annealed ingot is held at 1180℃ for 5 hours, then removed from the furnace and upsetting is performed with a downsetting amount of 20% and a downsetting time of 40 seconds to obtain a primary intermediate billet with a length of 0.8L;

[0122] ② The primary intermediate billet is held at 1180℃ for 5 hours, then taken out of the furnace and drawn out for 90 seconds to obtain a secondary intermediate billet of length L.

[0123] ③ The secondary intermediate billet is kept at 1160℃ for 4 hours, then out of the furnace for upsetting treatment, with a pressing amount of 30% and a pressing time of 45 seconds, to obtain a tertiary intermediate billet with a length of 0.7L;

[0124] ④ The three intermediate billets are held at 1160℃ for 4 hours, then taken out of the furnace and stretched for 100 seconds to obtain a four intermediate billet with a length of 1.2L.

[0125] ⑤ The four intermediate billets are held at 1140℃ for 4 hours, then out of the furnace for upsetting, with a pressing amount of 40% and a pressing time of 50 seconds, to obtain a five intermediate billet with a length of 0.6L.

[0126] ⑥ The fifth intermediate billet is held at 1140℃ for 4 hours, then taken out of the furnace and drawn out for 120 seconds to obtain a sixth intermediate billet with a size of Φ235mm±10mm.

[0127] ⑦ Preheat the hammer of the radial forging machine to 500℃ for 4 hours; keep the six intermediate billets at 1130℃ for 5 hours, and forge them on the radial forging machine to a size of Φ165mm±10mm. Air cool to room temperature, and machine to remove the black skin from the surface to a finished bar with a specification of Φ150mm.

[0128] (4) Heat treatment of the finished bar stock; the heat treatment includes: heat treatment at 1150℃ for 4 hours, air cooling to room temperature; then heat treatment at 860℃ for 6 hours and air cooling, then heat treatment at 780℃ for 24 hours and air cooling to room temperature.

[0129] Example 2

[0130] This embodiment provides a nickel-based wrought superalloy comprising the following components by mass percentage: Cr 11.0%, Co 15.5%, W 3.5%, Mo 4.6%, Al 3.8%, Ti 2.7%, Nb 3.4%, C 0.05%, B 0.006%, V 0.5%, Zr 0.06%, Si 0.02%, Re 0.02%, Sc 0.03%, La 0.02%, Ce 0.03%, Nd 0.04%, Mg 0.04%, P 0.012%, S 0.0006%, and Ni balance.

[0131] The preparation method of nickel-based wrought superalloys includes the following steps:

[0132] (1) Prepare a consumable ingot with a specification of Φ508mm by a triple smelting process of vacuum induction melting + protective atmosphere electroslag remelting + vacuum consumable remelting according to the alloy composition. Polish the surface of the consumable ingot and spray it with conventional anti-oxidation coating.

[0133] During vacuum induction melting, rare earth elements Sc, La, Ce, and Nd are added in the form of intermediate alloys Ni-Sc, Ni-La, Ni-Ce, and Ni-Nd respectively during the low-temperature refining stage at 1480–1510℃; the full melting temperature is controlled at 1540–1550℃, the refining temperature is controlled at 1540–1550℃, and the casting is carried out at 1440–1450℃.

[0134] During electroslag remelting under a protective atmosphere, the argon flow rate is 150±10 L / min;

[0135] During vacuum arc remelting, the vacuum arc remelting rate is 3.2 kg / min.

[0136] (2) The self-consumable ingot obtained in step (1) is subjected to multi-stage homogenization annealing to obtain an annealed ingot with a length of L. The multi-stage homogenization annealing includes: heating to 500℃ at 90℃ / h and holding for 4h, then heating to 1000℃ at 80℃ / h and holding for 8h, then heating to 1130℃ at 50℃ / h and holding for 20h, then heating to 1170℃ at 10℃ / h and holding for 30h, then heating to 1200℃ at 5℃ / h and holding for 35h, and then furnace cooling.

[0137] (3) The annealed ingot obtained in step (2) is subjected to upsetting and forging, specifically including the following steps:

[0138] ① The annealed ingot is held at 1170℃ for 5 hours, then removed from the furnace and upsetting is performed with a downsetting amount of 20% and a downsetting time of 40 seconds to obtain a primary intermediate billet with a length of 0.8L;

[0139] ② The primary intermediate billet is held at 1170℃ for 5 hours, then taken out of the furnace and drawn out for 90 seconds to obtain a secondary intermediate billet of length L.

[0140] ③ The secondary intermediate billet is kept at 1150℃ for 4 hours, then out of the furnace for upsetting treatment, with a pressing amount of 30% and a pressing time of 45 seconds, to obtain a tertiary intermediate billet with a length of 0.7L;

[0141] ④ The three intermediate billets are held at 1150℃ for 4 hours, then taken out of the furnace and stretched for 100 seconds to obtain a four intermediate billet with a length of 1.2L.

[0142] ⑤ The four intermediate billets are held at 1130℃ for 4 hours, then out of the furnace for upsetting, with a pressing amount of 40% and a pressing time of 50 seconds, to obtain a five intermediate billet with a length of 0.6L.

[0143] ⑥ The fifth intermediate billet is kept at 1130℃ for 4 hours, then taken out of the furnace and drawn out for 120 seconds to obtain a sixth intermediate billet with a size of Φ320mm±10mm.

[0144] ⑦ Preheat the hammer of the radial forging machine to 500℃ for 4 hours; keep the six intermediate billets at 1120℃ for 5 hours, and forge them on the radial forging machine to a size of Φ265mm±10mm. Air cool to room temperature, and machine to remove the black skin from the surface to a finished bar with a specification of Φ250mm.

[0145] (4) Heat treatment of the finished bar stock; the heat treatment includes: heat treatment at 1150℃ for 4 hours, air cooling to room temperature; then heat treatment at 860℃ for 6 hours and air cooling, then heat treatment at 780℃ for 24 hours and air cooling to room temperature.

[0146] Example 3

[0147] This embodiment provides a nickel-based wrought superalloy comprising the following components by mass percentage: Cr 11.5%, Co 15.5%, W 3.2%, Mo 4.4%, Al 3.7%, Ti 2.8%, Nb 3.5%, C 0.04%, B 0.005%, V 0.7%, Zr 0.05%, Si 0.03%, Re 0.02%, Sc 0.04%, La 0.03%, Ce 0.02%, Nd 0.03%, Mg 0.03%, P 0.013%, S 0.0004%, and Ni balance.

[0148] The preparation method of nickel-based wrought superalloys includes the following steps:

[0149] (1) Prepare a consumable ingot with a specification of Φ508mm by a triple smelting process of vacuum induction melting + protective atmosphere electroslag remelting + vacuum consumable remelting according to the alloy composition. Polish the surface of the consumable ingot and spray it with conventional anti-oxidation coating.

[0150] During vacuum induction melting, rare earth elements Sc, La, Ce, and Nd are added in the form of intermediate alloys Ni-Sc, Ni-La, Ni-Ce, and Ni-Nd respectively during the low-temperature refining stage at 1480–1510℃; the full melting temperature is controlled at 1540–1550℃, the refining temperature is controlled at 1540–1550℃, and the casting is carried out at 1440–1450℃.

[0151] During electroslag remelting under a protective atmosphere, the argon flow rate is 150±10 L / min;

[0152] During vacuum arc remelting, the vacuum arc remelting rate is 3.2 kg / min.

[0153] (2) The consumable ingot obtained in step (1) is subjected to multi-stage homogenization annealing to obtain an annealed ingot with a length of L. The multi-stage homogenization annealing includes: heating at 90℃ / h to 500℃ and holding for 4h, then heating at 80℃ / h to 1000℃ and holding for 8h, then heating at 50℃ / h to 1120℃ and holding for 25h, then heating at 10℃ / h to 1160℃ and holding for 25h, then heating at 5℃ / h to 1190℃ and holding for 38h, and then furnace cooling.

[0154] (3) The annealed ingot obtained in step (2) is subjected to upsetting and forging, specifically including the following steps:

[0155] ① The annealed ingot is held at 1180℃ for 5 hours, then removed from the furnace and upsetting is performed with a downsetting amount of 20% and a downsetting time of 40 seconds to obtain a primary intermediate billet with a length of 0.8L;

[0156] ② The primary intermediate billet is held at 1180℃ for 4 hours, then taken out of the furnace and stretched for 90 seconds to obtain a secondary intermediate billet of length L.

[0157] ③ The secondary intermediate billet is kept at 1160℃ for 4 hours, then out of the furnace for upsetting treatment, with a pressing amount of 30% and a pressing time of 45 seconds, to obtain a tertiary intermediate billet with a length of 0.7L;

[0158] ④ The three intermediate billets are held at 1160℃ for 4 hours, then taken out of the furnace and stretched for 100 seconds to obtain a four intermediate billet with a length of 1.2L.

[0159] ⑤ The four intermediate billets are held at 1140℃ for 4 hours, then out of the furnace for upsetting, with a pressing amount of 40% and a pressing time of 50 seconds, to obtain a five intermediate billet with a length of 0.6L.

[0160] ⑥ The fifth intermediate billet is kept at 1140℃ for 4 hours, then taken out of the furnace and drawn out for 120 seconds to obtain a sixth intermediate billet with a size of Φ320mm±10mm.

[0161] ⑦ The sixth intermediate billet is kept at 1130℃ for 4 hours, then taken out of the furnace and drawn out for 90 seconds to obtain a seventh intermediate billet with a size of Φ280mm±10mm.

[0162] ⑧ The hammer of the radial forging machine is preheated at 500℃ for 4 hours. The six intermediate billets are held at 1110℃ for 5 hours. After being taken out of the furnace, they are forged on the radial forging machine to a size of Φ230mm±10mm. They are then air-cooled to room temperature and machined to remove the black skin from the surface to a finished bar with a specification of Φ200mm.

[0163] (4) Heat treatment of the finished bar stock; the heat treatment includes: heat treatment at 1150℃ for 4 hours, air cooling to room temperature; then heat treatment at 860℃ for 6 hours and air cooling, then heat treatment at 780℃ for 24 hours and air cooling to room temperature.

[0164] Example 4

[0165] This embodiment refers to the nickel-based wrought superalloy and its preparation method in Embodiment 1, the only difference being the alloy composition.

[0166] The nickel-based wrought superalloy of this embodiment comprises the following components by mass percentage: Cr 10.5%, Co 15.1%, W 3.4%, Mo 4.5%, Al 4.1%, Ti 2.6%, Nb 3.3%, C 0.06%, B 0.005%, V 0.6%, Zr 0.05%, Si 0.015%, Re 0.02%, Sc 0.06%, La 0.02%, Ce 0.02%, Nd 0.05%, Mg 0.03%, P 0.012%, S 0.0007%, and Ni balance.

[0167] Example 5

[0168] This embodiment refers to the nickel-based wrought superalloy and its preparation method in Embodiment 1, the only difference being the alloy composition.

[0169] The nickel-based wrought superalloy of this embodiment comprises the following components by mass percentage: Cr 10.5%, Co 15.0%, W 3.5%, Mo 4.4%, Al 3.9%, Ti 2.8%, Nb 3.3%, C 0.06%, B 0.005%, V 0.6%, Zr 0.05%, Si 0.015%, Re 0.04%, Sc 0.06%, La 0.02%, Ce 0.02%, Nd 0.05%, Mg 0.03%, P 0.012%, S 0.0005%, and Ni balance.

[0170] Example 6

[0171] This embodiment refers to the nickel-based wrought superalloy and its preparation method in Embodiment 1, the only difference being the alloy composition.

[0172] The nickel-based wrought superalloy of this embodiment comprises the following components by mass percentage: Cr 10.5%, Co 15.0%, W 3.4%, Mo 4.5%, Al 3.9%, Ti 2.8%, Nb 3.3%, C 0.06%, B 0.005%, V 0.6%, Zr 0.05%, Si 0.015%, Re 0.03%, Sc 0.08%, La 0.08%, Ce 0.05%, Nd 0.05%, Mg 0.03%, P 0.015%, S 0.0006%, and Ni balance.

[0173] Example 7

[0174] Example 7 is a nickel-based wrought superalloy and its preparation method as described in Example 1. The only difference is that the homogenization annealing in step (2) of the preparation method is different.

[0175] The homogenization annealing in step (2) of this embodiment includes: heating to 500°C at 90°C / h and holding for 4 hours, then heating to 1000°C at 80°C / h and holding for 8 hours, then heating to 1150°C at 50°C / h and holding for 15 hours, then heating to 1200°C at 5°C / h and holding for 25 hours, and then furnace cooling.

[0176] Example 8

[0177] Example 8 is a nickel-based wrought superalloy and its preparation method as described in Example 1, with the only difference being that step (3) is different in the preparation method.

[0178] Step (3) of this embodiment includes: upsetting and forging the annealed ingot obtained in step (2), specifically including the following steps:

[0179] ① The annealed ingot is held at 1180℃ for 5 hours, then removed from the furnace and upsetting is performed with a downsetting amount of 20% and a downsetting time of 40 seconds to obtain a primary intermediate billet with a length of 0.8L;

[0180] ② The primary intermediate billet is held at 1180℃ for 5 hours, then taken out of the furnace and drawn out for 90 seconds to obtain a secondary intermediate billet of length L.

[0181] ③ The secondary intermediate billet is kept at 1180℃ for 4 hours, then removed from the furnace and upsetting is performed with a pressing amount of 30% and a pressing time of 45 seconds to obtain a tertiary intermediate billet with a length of 0.7L.

[0182] ④ The three intermediate billets are held at 1180℃ for 4 hours, then taken out of the furnace and stretched for 100 seconds to obtain a four intermediate billet with a length of 1.2L.

[0183] ⑤ The four intermediate billets are held at 1180℃ for 4 hours, then out of the furnace for upsetting, with a pressing amount of 40% and a pressing time of 50 seconds, to obtain a five intermediate billet with a length of 0.6L.

[0184] ⑥ The fifth intermediate billet is kept at 1180℃ for 4 hours, then taken out of the furnace and drawn out for 120 seconds to obtain a sixth intermediate billet with a size of Φ235mm±10mm.

[0185] ⑦ Preheat the hammer of the radial forging machine to 500℃ for 4 hours; keep the six intermediate billets at 1130℃ for 5 hours, and forge them on the radial forging machine to a size of Φ165mm±10mm. Air cool to room temperature, and machine to remove the black skin from the surface to a finished bar with a specification of Φ150mm.

[0186] Comparative Example 1

[0187] Comparative Example 1 refers to the nickel-based wrought superalloy and its preparation method in Example 1, the difference being that the alloy composition is different.

[0188] The nickel-based wrought superalloy of Comparative Example 1 comprises the following components by mass percentage: Cr 10.5%, Co 15.0%, W 3.4%, Mo 4.5%, Al 3.9%, Ti 2.8%, Nb 3.3%, C 0.06%, B 0.005%, V 0.6%, Zr 0.05%, P 0.015%, S 0.0006%, and Ni balance.

[0189] Comparative Example 2

[0190] Comparative Example 2 refers to the nickel-based wrought superalloy and its preparation method in Example 1, the difference being that the alloy composition is different.

[0191] The nickel-based wrought superalloy of Comparative Example 2 comprises the following components by mass percentage: Cr 10.5%, Co 15.0%, W 3.4%, Mo 4.5%, Al 3.9%, Ti 2.8%, Nb 4.1%, C 0.06%, B 0.005%, V 0.6%, Zr 0.05%, Si 0.015%, Re 0.03%, Sc 0.06%, La 0.02%, Ce 0.02%, Nd 0.05%, Mg 0.03%, P 0.012%, S 0.0006%, and Ni balance.

[0192] Comparative Example 3

[0193] Comparative Example 3 refers to the nickel-based wrought superalloy and its preparation method in Example 1. The difference is that the homogenization annealing in step (2) of the preparation method is different.

[0194] The homogenization annealing in step (2) of Comparative Example 3 includes: heating to 1210°C at 5°C / h and holding for 30h, followed by furnace cooling.

[0195] Comparative Example 4

[0196] Comparative Example 4 refers to the nickel-based wrought superalloy and its preparation method in Example 2, the difference being that in the preparation method, step (3) does not include radial forging.

[0197] Step (3) of Comparative Example 4 includes: preparing six intermediate billets according to steps (3) ① to ⑥ of Example 2; then keeping the six intermediate billets at 1120°C for 5 hours, drawing them after they are taken out of the furnace for 120 seconds, drawing them and rounding them to Φ280mm±10mm, air cooling them to room temperature, and machining them to remove the black skin from the surface to a finished bar with a specification of Φ250mm.

[0198] Experimental Example

[0199] The bar samples prepared in each embodiment and comparative example were subjected to metallographic observation and mechanical property testing. The grain structure is shown in Table 1, and the mechanical properties are shown in Table 2. Figure 3 and Figure 4 The images show the grain structure of the rods obtained in Example 1 and Comparative Example 1 of this invention, respectively. Figure 5 The microstructure of the annealed alloy ingot prepared in Comparative Example 1 is shown in the diagram. Figure 6 and Figure 7 The images show the tensile fracture morphology of the bar obtained in Example 1 of this invention at 800℃ and the fracture morphology of the long-term specimen at 800℃ / 500MPa. Figure 8 and Figure 9 The images show the tensile fracture morphology of the bar obtained in Comparative Example 1 at 800℃ and the fracture morphology of the long-term specimen at 800℃ / 500MPa.

[0200] Table 1 Metallographic structure of different bars

[0201]

[0202]

[0203] Table 2. Test results of mechanical properties of different bars

[0204]

[0205] The test results above show that the nickel-based wrought superalloy of the present invention has high homogeneity and low segregation characteristics. The average grain size of the fine-grained rod is finer than grade 8, and the grade difference is less than 2 grades. The room temperature tensile strength is ≥1655MPa, the 750℃ tensile strength is ≥1260MPa, the 750℃ / 620MPa creep rupture life is ≥145h, and the 800℃ / 500MPa creep rupture life is ≥72h.

[0206] Finally, it should be noted that the above embodiments are only used to illustrate the technical solutions of the present invention, and not to limit them; although the present invention has been described in detail with reference to the foregoing embodiments, those skilled in the art should understand that modifications can still be made to the technical solutions described in the foregoing embodiments, or equivalent substitutions can be made to some or all of the technical features; and these modifications or substitutions do not cause the essence of the corresponding technical solutions to deviate from the scope of the technical solutions of the embodiments of the present invention.

Claims

1. A high-strength, high-toughness nickel-based high-temperature alloy resistant to deformation at 800℃, characterized in that, Includes the following components by mass percentage: Cr 10%–12%, Co 14%–16%, W 2.5%–3.5%, Mo 3.75%–5%, Al 3.5%–4.5%, Ti 2.5%–3.1%, Nb 3%–3.8%, C 0.01%–0.08%, B 0.005%–0.05%, V 0.01%–1%, Zr 0.025%–0.075%, Si 0.005%–0.3%, Re 0.02%–0.05%, Sc 0.001%–0.08%, La 0.001%–0.08%, Ce 0.001%–0.05%, Nd 0.001%–0.05%, Mg 0.001%–0.15%, with the balance being Ni and unavoidable impurities; The mass ratio of Al to Ti is 1.15 to 1.

45. The mass ratio of Mo to W, Mo / W, is 1.3 to 1.45; The sum of the mass percentages of Sc, La, Ce, and Nd satisfies: Sc + La + Ce + Nd ≤ 0.2%; The preparation method of the high-strength and high-toughness nickel-based high-temperature alloy resistant to 800℃ includes the following steps: (a) According to the alloy composition, alloy ingots are obtained through a three-stage smelting process; (b) The alloy ingot is subjected to multi-stage homogenization annealing at 1100-1220°C to obtain an annealed ingot; (c) The annealed ingot is upset at 1060-1180°C and then radially forged at 1020-1150°C; The alloy after radial forging is subjected to heat treatment, which includes solution treatment and aging treatment; the solution treatment includes holding at 1100-1150℃ for 2-6 hours; the aging treatment includes holding at 800-900℃ for 2-6 hours followed by air cooling, and then holding at 750-820℃ for 8-24 hours followed by air cooling. The multi-stage homogenization annealing includes: holding the alloy ingot at 490-510℃ for more than 4 hours, then raising the temperature to 990-1010℃ and holding it for more than 8 hours, then raising the temperature to 1110-1140℃ and holding it for more than 15 hours, then raising the temperature to 1150-1180℃ and holding it for more than 25 hours, then raising the temperature to 1190-1220℃ and holding it for more than 30 hours, and then furnace cooling. The upsetting and drawing of the billet is carried out using a high-speed forging mill, with an upsetting deformation of 10% to 55% per pass and a drawing deformation of 15% to 60% per pass. In the upsetting and drawing process, the forging temperature is gradually reduced in stages within the range of 1060 to 1180°C.

2. The high-strength, high-toughness nickel-based high-temperature alloy resistant to 800℃ deformation according to claim 1, characterized in that, The sum of the mass percentages of Al, Ti, and Nb satisfies: Al + Ti + Nb ≥ 9.5%.

3. The high-strength, high-toughness nickel-based high-temperature alloy resistant to deformation at 800℃ according to claim 1, characterized in that, The sum of the mass percentages of Mo and W satisfies: Mo + W ≥ 7.5%.

4. The high-strength, high-toughness nickel-based high-temperature alloy resistant to deformation at 800℃ according to claim 1, characterized in that, The sum of the mass percentages of C, B, and Zr satisfies: C + B + Zr ≤ 0.15%.

5. The high-strength, high-toughness nickel-based high-temperature alloy resistant to deformation at 800℃ according to claim 1, characterized in that, In the radial forging process, an alloy of a predetermined specification is prepared by a multi-pass deformation method in a single heat.

6. The high-strength, high-toughness nickel-based high-temperature alloy resistant to deformation at 800℃ according to claim 1, characterized in that, In the radial forging process, the holding time at 1020–1150℃ is 30–300 min.

7. The high-strength, high-toughness nickel-based high-temperature alloy resistant to deformation at 800℃ according to claim 1, characterized in that, The heat-treated alloy has at least one of the following characteristics: (1) The average grain size is grade 8 or above, and the grade difference is < 2 grades; (2) The content of γ' phase is 50wt%~58wt%; (3) Tensile strength at room temperature ≥ 1510 MPa, yield strength at room temperature ≥ 1130 MPa; (4) Tensile strength at 750℃ ≥ 1110 MPa, yield strength at 750℃ ≥ 960 MPa; (5) Tensile strength at 800℃ ≥ 1000 MPa, yield strength at 800℃ ≥ 900 MPa; (6) Duration life at 750℃ / 620MPa ≥60h; (7) Duration life at 800℃ / 500MPa ≥45h.

8. A turbine disk, characterized in that, It is prepared using the high-strength and high-toughness nickel-based high-temperature alloy that can withstand 800℃ deformation as described in any one of claims 1 to 7.

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