800°c-resistant high-strength and high-toughness nickel-based wrought superalloy and preparation method therefor, and turbine disc

Abstract

The present application relates to the technical field of superalloy, in particular to a high-strength and high-toughness nickel-based wrought superalloy resistant to 800°C, preparation method thereof and turbine disk. The nickel-based wrought superalloy comprises 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%, and the rest is Ni. The present application reasonably and coordinately designs the components of the nickel-based high-temperature alloy, coarsens the γ' phase through process regulation and control, which makes the alloy achieve good hot working performance in engineering, while obtaining the turbine disc meeting the requirement through heat treatment.

Description

TECHNICAL FIELD

[0001] The present application relates to the technical field of superalloy, in particular to high-strength and high-toughness nickel-based wrought superalloy resistant to 800°C, preparation method thereof and turbine disk.BACKGROUND ART

[0002] The turbine disk is an important hot end component of an aircraft engine, has a harsh working environment, and is mainly made of wrought superalloys. With the development of the aviation industry, the outlet temperature of aircraft engines has gradually increased, so the temperature-bearing capacity of nickel-based wrought superalloys used for turbine disks has also increased, which requires the wrought superalloys to have excellent medium-temperature (600-800°C) service performance. The main methods of strengthening and toughening design for the wrought superalloys are: addition of of Al / Ti and Nb to form γ' or γ" phase for precipitation strengthening, W and Mo for solid solution strengthening, and micro-alloying elements such as B, Zr and C for grain boundary strengthening.

[0003] Generally, there are two main process routes for highly alloyed superalloys: the powder metallurgy route reduces the severity of segregation, but it is difficult to solve defects such as prior particle boundary, inclusions and heat-induced holes, which seriously affects the mechanical properties of powder metallurgy superalloys such as low-cycle fatigue; the forging process route is the current mainstream process, which is mature and suitable for industrial application.

[0004] At present, the main typical grades of the nickel-based wrought superalloys are GH4169 alloy resistant to 650°C, GH169D alloy resistant to 700°C, GH4720Li alloy resistant to 750°C and GH4065A alloy. These alloys have gradually increased alloying degree, and a high alloying degree will significantly reduce the plasticity of the alloy during hot deformation process and in service (medium temperature) environments; the higher the alloying degree, the more serious the segregation of elements in the high-temperature alloy ingots and the more complex the solidification segregation phase. For large-sized alloy ingots, the segregation of elements is serious and the inheritance of the segregated organization is difficult to eliminate, so the ingot specifications are limited to small ingots. But the problems such as narrow hot working window, large deformation resistance, sensitivity to cracking make the microstructure control of the cogging process extremely difficult, the yield rate extremely low, and the microstructure uneven, which has become a key problem restricting the preparation and application of high-quality bars of the highly alloyed superalloys.

[0005] In view of the above analysis, the present application is proposed.SUMMARY OF THE PRESENT APPLICATION

[0006] One purpose of the present application is to provide a high-strength and high-toughness nickel-based wrought superalloy resistant to 800°C, which not only has both strength and toughness at service temperature, but also has applicable machinability and can provide the necessary high temperature strength nickel-based wrought superalloys for high thrust-to-weight ratio engines.

[0007] Another purpose of the present application is to provide a preparation method of the high-strength and high-toughness nickel-based wrought superalloy resistant to 800°C.

[0008] Another purpose of the present application is to provide a turbine disk, which is prepared by using the above high-strength and high-toughness nickel-based wrought superalloy resistant to 800°C.

[0009] In order to achieve the above-mentioned purposes of the present application, on the one hand, the present application provides a high-strength and high-toughness nickel-based wrought superalloy resistant to 800°C, which comprises 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%, and the rest is Ni and unavoidable impurities.

[0010] In specific embodiments of the present application, the sum of the mass percentages of Al, Ti and Nb satisfies: Al+Ti+Nb≥9.5%; and the mass ratio Al / Ti of Al to Ti is 1.15-1.45.

[0011] In specific embodiments of the present application, the sum of the mass percentages of Mo and W satisfies: Mo+W≥7.5%; and the mass ratio Mo / W of Mo to W is 1.3-1.45.

[0012] In specific embodiments of the present application, the sum of the mass percentages of C, B and Zr satisfies: C+B+Zr≤0.15%.

[0013] In specific embodiments of the present application, the sum of the mass percentages of Sc, La, Ce and Nd satisfies: Sc+La+Ce+Nd≤0.2%.

[0014] In specific embodiments of the present application, the unavoidable impurities include any one or more of S, P, N, O, Cu, Ag, Si, Mn, Ca, Sn and Pb. Further, by mass percentage, the impurity elements in the alloy 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%, Pb≤0.001%.

[0015] On the other hand, the present application provides a method for preparing any one of the above-mentioned high-strength and high-toughness nickel-based wrought superalloys resistant to 800°C, and the method comprises the following steps: (a) preparing an alloy ingot by a triple smelting process (VIM +ESR +VAR) according to alloy components; (b) subjecting the alloy ingot to multi-stage homogenization at 1100-1220° C to obtain an annealed ingot; (c) subjecting the annealed ingot to upsetting and drawing at 1060-1180°C, and then subjecting to radial forging at 1020-1150°C.

[0016] In specific embodiments of the present application, the multi-stage homogenization includes: keeping the alloy ingot at 490-510°C for more than 4 h, then heating to 990-1010°C and keeping for more than 8 h, then heating to 1110-1140°C and keeping for more than 15 h, then heating to 1150-1180°C and keeping for more than 25 h, then heating to 1190-1220°C and keeping for more than 30 h, and then furnace cooling.

[0017] In specific embodiments of the present application, adopting a high-speed forging press to perform the upsetting and drawing , and each of the upsetting deformation amount is 10%~55%, and each of the drawing deformation amount is 15%~60%. Furthermore, a sequential heat temperature decrease from 1060°C to 1180°C was implemented in the process of the upsetting and drawing..

[0018] In specific embodiments of the present application, in the upsetting and drawing the post-deformation holding times is 1~6 h.

[0019] In specific embodiments of the present application, in the radial forging , multi-pass forging in one heat is used to prepare an alloy of preset specifications. Furthermore, in the radial forging, the holding time at 1020°C-1150°C is 30-300 min.

[0020] In specific embodiments of the present application, the method further includes: performing heat treatment including solid solution treatment and aging of the alloy after the radial forging. Further, the solid solution treatment includes: holding at 1100-1150°C for 2-6 h. The aging includes holding at 800-900°C for 2-6 h then air cooled and holding at 750-820°C for 8-24 h then air cooled.

[0021] In specific embodiments of the present application, the average grain size of the heat-treated alloy is Grade 8 or finer, and the grain size variation is less than 2 grades. Furthermore, the content of γ' phase of the heat-treated alloy is 50wt%~58wt%.

[0022] In specific embodiments of the present application, the tensile properties of the heat-treated alloy satisfy: The tensile strength at room temperature is ≥1510MPa, the yield strength at room temperature is ≥1130MPa.

[0023] The tensile strength at 750°C is ≥1110MPa, the yield strength at 750°C is ≥960MPa.

[0024] The tensile strength at 800°C is ≥1000MPa, and the yield strength at 800°C is ≥900MPa.

[0025] In specific embodiments of the present application, the creep rupture life of the heat-treated alloy satisfies:

[0026] The creep rupture life of 60h minimum at 750°C under 620MPa.

[0027] The creep rupture life of 45h minimum at 800°C under 500MPa.

[0028] Another aspect of the present application provides a turbine disk, which is made of any one of the above-mentioned high-strength and high-toughness nickel-based wrought superalloys resistant to 800°C.

[0029] Compared with the prior art, the present application has the following beneficial effects: (1) The present application achieves high strength-toughness and excellent processability in nickel-based superalloys through rational compositional design incorporating specific rare earth elements and trace elements such as C, B, and Zr. This composition enhances grain boundary cohesion, reduces or eliminates brittleness tendencies within the service temperature range, and enables processing manufacturability by controlling the coarsening of γ' phase during fabrication. The optimized microstructure prevents medium-temperature embrittlement and low plasticity at operational temperatures, ensuring the alloy simultaneously exhibits exceptional mechanical properties and process adaptability. (2) The present application adopts a highly homogeneous and low-segregation alloy through compositional optimization and process control, employing multi-stage homogenization annealing, combined upsetting and drawing, and radial forging. This approach ensures high strength-toughness at 800°C. By tailoring the process parameters, the strengthening effect of the γ' phase is weakened, refining the microstructure into a fine-grained structure to achieve good machinability. (3) The nickel-based wrought superalloy of the present application meets the high thermal strength material requirements of high thrust-to-weight ratio engines, which is of great significance to the development of the aviation field. A BRIEF DESCRIPTION OF DRAWINGS

[0030] In order to illustrate the specific embodiments of the present application or the technical solutions in the prior art more clearly, the drawings required for using in the specific embodiments or the description of the prior art will be briefly introduced below. Obviously, the drawings described below correspond to some embodiments of the present application, for ordinary technicians in this field, other drawings can be obtained based on these drawings without paying creative work. Fig.1 is a diagram showing the non-equilibrium solidification segregation characteristics such as alloy segregation phase and element segregation of an alloy system according to an embodiment of the present application. Fig.2 is a multi-stage high-temperature diffusion homogenization annealing temperature curve provided by an embodiment of the present application. Fig.3 is a grain structure diagram of the bar material obtained in Embodiment 1 of the present application. Fig.4 is a grain structure diagram of the bar obtained in Comparative Embodiment 1. Fig.5 is a microstructure diagram of the annealed alloy ingot prepared in Comparative Embodiment 1. Fig.6 is a tensile fracture morphology at 800°C of the bar obtained in Embodiment 1 of the present application. Fig.7 is a fracture morphology of long-term specimen under an 800°C / 500MPa environment of the bar prepared in Embodiment 1 of the present application. Fig.8 is a tensile fracture morphology at 800°C of the bar prepared in Comparative Embodiment 1. Fig. 9 is a fracture morphology of long-term specimen under an 800°C / 500MPa environment of the bar prepared in Comparative Embodiment 1 of the present application. DETAILED DESCRIPTION OF THE PRESENT APPLICATION

[0031] The technical scheme of the present application will be clearly and completely described below in conjunction with the accompanying drawings and specific embodiments, but it will be understood by those skilled in the art that the following described embodiments are part of embodiments of the present application rather than all embodiments, and are only used to illustrate the present application but should not be considered as limiting the scope of the present application. Based on the embodiments of the present application, all other embodiments obtained by those skilled in the art without making creative work belong to the scope of protection of the present application. Carrying out the specific conditions without being specified in the embodiments according to normal conditions or conditions recommended by the manufacturer. The reagents or instruments used without the manufacturer being specified are all conventional products that can be purchased commercially.

[0032] On the one hand, the present application provides a high-strength and high-toughness nickel-based wrought superalloy resistant to 800°C, which comprises 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%, and the rest is Ni and unavoidable impurities.

[0033] The present application improves the hot working performance and mechanical properties of the alloy by performing rational and coordinated design of the components of the nickel-based high-temperature alloy. The alloy of the present application belongs to a high-alloyed hard-to-deform nickel-based high-temperature alloy with the γ' phase content in the alloy about 10% higher than that in the 750°C-resistant alloy, at the same time, the content of solid solution strengthening elements in the alloy is also limited to a high level, which make the alloy have excellent high-temperature strength at 800°C. The addition of rare earth elements such as Sc and grain boundary strengthening elements such as Zr into the alloy improves the grain boundary bonding strength, makes the equal strength transformation temperature between the grain boundary and the interior of the grain transfer to a higher temperature to make the alloy obtain better toughness, and avoids medium-temperature brittleness at around 800°C. In conjunction with the process control during the hot working process, weakening the strengthening effect of the γ' phase and regulating the alloy to have a fine-grained structure, which can make the alloy be prepared by hot working in engineering.

[0034] A large amount of solid solution strengthening elements, precipitation strengthening elements, grain boundary strengthening elements and rare earth elements are added into the alloy of the present application and have the specific composition characteristics as follows.

[0035] Adding the precipitation strengthening elements of Al, Ti and Nb into the alloy to make the γ' phase content of 50wt%~58wt%, which is significantly higher than other high-heat strength nickel-based wrought superalloys. A large amount of precipitation strengthening phases ensure that the alloy has sufficient strength. The Nb element enters the γ' phase to form Ni 3 (Al, Ti, Nb), which can increase the volume fraction and the reverse domain boundary energy of the γ' phase in the alloy, thereby improving the strength of the alloy. In order to take into account the strength and processability of the alloy, the present application further preferably regulates the sum of the mass percentages of Al, Ti and Nb to satisfies Al+Ti+Nb ≥ 9.5%, and the mass ratio Al / Ti of Al and Ti is 1.15~1.45.

[0036] For example, in different embodiments, the mass percentage of Al in the present application can be 3.5%, 3.6%, 3.7%, 3.8%, 3.9%, 4%, 4.1%, 4.2%, 4.3%, 4.4%, 4.5% or a range consisting of any two of them; the mass percentage of Ti can be 2.5%, 2.6%, 2.%, 2.8%, 2.9%, 3%, 3.1% or a range consisting of any two of them; 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%~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 a range consisting of any two thereof; the mass ratio Al / Ti of Al and Ti can be 1.15, 1.2, 1.25, 1.3, 1.35, 1.4, 1.45 or a range consisting of any two thereof.

[0037] Adding solid solution strengthening elements W and Mo into the alloy to cause significant lattice distortion and reduce stacking fault energy, thereby improving the yield strength and creep properties of the alloy. However, when excessive or improperly proportioned W and Mo are added, a harmful brittle phase TCP phase will be generated during the long-term use of the alloy, which makes cracks initiate and expand at the interface between the TCP phase and the matrix under the action of thermal-mechanical coupling, thereby resulting in deterioration of the alloy performance. Therefore, the present application further regulates the sum of the mass percentages of Mo and W to satisfy: W+Mo≥7.5%, and the mass ratio Mo / W of M to W is 1.3~1.45, so that the alloy has a sufficient solid solution strengthening effect.

[0038] For example, in different embodiments, the mass percentage of W in the present application can be 2.5%, 2.6%, 2.7%, 2.8%, 2.9%, 3%, 3.1%, 3.2%, 3.3%, 3.4%, 3.5% or a range of any two thereof; 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%, 5%. Furthermore, the sum of the mass percentages of Mo and W is 7.5%~8.2%, such as 7.5%, 7.6%, 7.7%, 7.8%, 7.9%, 8%, 8.1%, 8.2% or a range consisting of any two thereof; the mass ratio Mo / W of Mo to W can be 1.3, 1.32, 1.35, 1.38, 1.4, 1.42, 1.45 or a range consisting of any two thereof.

[0039] The C element added into the alloy of the present application segregates to the grain boundary or forms grain boundary carbides, and precipitates in the form of blocks or particles at the grain boundary, which can inhibit the growth of abnormal grain, prevent grain boundary sliding, strengthen the grain boundary, improve the lasting strength and plasticity. However, the excessive addition of C element causes the carbides to precipitate continuously along the grain boundary in the form of a thin film, which makes the grain boundary embrittled and the alloy exhibits notch sensitivity. Therefore, the C content in the present application is controlled in the range of 0.01%~0.08%. For example, 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 a range consisting of any two of them.

[0040] The elements of Zr and B added into the alloy of the present application segregate to the grain boundaries, which improves the grain boundary bonding force, reduces the interface energy, improves the grain boundary morphology, improves the creep endurance strength and the low-cycle fatigue performance. At the same time, the Zr element can also be used as a purifier and react with elements such as C and S in the alloy to generate carbon sulfides or carbides, thereby removing harmful impurity elements such as S and P. while adding excessive B elements will cause the precipitation of boride at the grain boundaries or induce the generation of hot cracks, therefore, in the present application, the Zr content is controlled at 0.025%~0.075%, and the B content is controlled at 0.005%~0.05%. For example, in different embodiments, the mass percentage of Zr can be 0.025%, 0.035%, 0.045%, 0.055%, 0.065%, 0.075% or a range consisting of any two thereof, and the mass percentage of B can be 0.005%, 0.008%, 0.01%, 0.02%, 0.03%, 0.04%, 0.05% or a range consisting of any two thereof.

[0041] The present application further preferably regulates the sum of the mass percentages of C, B and Zr to satisfy: C+B+Zr≤0.15%, so as to improve the grain boundary bonding force, improve the endurance strength and plasticity, reduce or eliminate the brittle tendency in the service temperature range, and avoid grain boundary embrittlement and the generation of thermal cracking, etc. For example, 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 a range consisting of any two of them.

[0042] The rare earth elements of Sc, La, Ce and Nd added into the alloy of the present application can reduce the critical supercooling degree of dendrite nucleation, generate component supercooling at the front of the solid-liquid interface in the solidification process of the alloy, reduce the critical nucleation radius and nucleation work, increase the nucleation rate, thereby refining the dendrite and effectively reducing segregation. At the same time, the rare earth elements can also remove O and purify the grain boundary, reduce or avoid the generation of intergranular cracks of the alloy in the medium temperature range, and at the same time, the rare earth elements can also improve the oxidation resistance of the alloy. But excessive rare earth elements are easily oxidized, and the rare earth-rich phase is not easy to be eliminated in the subsequent homogenization and thermal deformation process. Therefore, the rare earth content in the present application should be strictly controlled, especially, controlling the Sc content at 0.001%~0.08%, controlling the La content at 0.001%~0.08%, controlling the Ce content at 0.001%~0.05%, and controlling the Nd content in the range of 0.001% ~0.05%. For example, in different embodiments, the mass percentage of Sc can be 0.001%, 0.005%, 0.01%, 0.02%, 0.04%, 0.06%, 0.08% or a range consisting of any two thereof, the mass percentage of La can be 0.001%, 0.005%, 0.01%, 0.02%, 0.04%, 0.06%, 0.08% or a range consisting of any two thereof, the mass percentage of Ce can be 0.001%, 0.005%, 0.01%, 0.02%, 0.03%, 0.04%, 0.05% or a range consisting of any two thereof, and the mass percentage of Nd can be 0.001%, 0.005%, 0.01%, 0.02%, 0.03%, 0.04%, 0.05% or a range consisting of any two thereof.

[0043] The sum of the mass percentages of Sc, La, Ce and Nd is further adjusted to satisfy Sc+La+Ce+Nd≤0.2%, such as 0.1%~0.2%, to optimize the above-mentioned properties. For example, in different embodiments, the sum of the mass percentages of Sc, La, Ce and Nd can be 0.1%, 0.11%, 0.12%, 0.13%, 0.14%, 0.15%, 0.16%, 0.17%, 0.18%, 0.19%, 0.2% or a range consisting of any two of them.

[0044] The Cr element in the alloy of the present application is the main element forming carbides and can also improve the oxidation resistance and corrosion resistance of the alloy. Too low Cr content will lead to a decrease in the oxidation resistance and corrosion resistance of the alloy and will lead to the formation of an α-phase, which is detrimental to the plasticity and toughness of the alloy. Therefore, the present application regulates the Cr content in the alloy to 10%~12%, which makes the alloy system of the present application take into account the improvement of oxidation resistance and corrosion resistance as well as plasticity and toughness by cooperating with the remaining elements. For example, in different embodiments, the mass percentage of Cr in the present application can be 10%, 10.2%, 10.5%, 10.8%, 11%, 11.2%, 11.5%, 11.8%, 12% or a range consisting of any two of them.

[0045] In specific embodiments of the present application, 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%, Pb≤0.001%.

[0046] In a specific embodiment of the present application, the high-strength and high-toughness nickel-based wrought superalloy resistant to 800 °C comprises the following components by mass percentage: Cr 10.5%~11.5%, Co 15%~15.5%, W 3.2%~3.5%, Mo 4.4%~4.6%, Al 3.7%~3.9%, Ti 2.7%~2.8%, Nb 3.3%~3.5%, C 0.04%~0.06%, B 0.005%~ 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%~0.04%, and the rest is Ni and inevitable impurities.

[0047] On the other hand, the present application provides a method for preparing any one of the above-mentioned the high-strength and high-toughness nickel-based wrought superalloys resistant to 800°C, and the method comprises the following steps: (a) preparing an alloy ingot by a triple smelting process according to alloy components; (b) subjecting the alloy ingot to multi-stage homogenization annealing at 1100-1220° C to obtain an annealed ingot; (c) subjecting the annealed ingot to upsetting and drawing cogging at 1060-1180°C, and then subjecting to radial forging at 1020-1150°C.

[0048] The material of the present application has a high alloying degree and requires high alloy purity. Through the triple smelting process of vacuum induction melting + electroslag remelting + vacuum induction remelting, combining with the alloy composition of the present application, an ingot with higher purity can be obtained.

[0049] In specific embodiments of the present application, adding the rare earth elements of Sc, La, Ce and Nd in the form of an intermediate alloy in the vacuum induction melting of the triple smelting; wherein the intermediate alloy is an Al-rare earth intermediate alloy or a Ni-rare earth intermediate alloy. For example, they can be added in the form of Al-Sc, Al-La, Al-Ce, Al-Nd, or in the form of Ni-Sc, Ni-La, Ni-Ce, Ni-Nd to improve the yield of the rare earth elements.

[0050] In specific embodiments of the present application, the intermediate alloy is added in the low temperature refining stage (1450-1580° C.) to improve the yield of rare earth elements.

[0051] In specific embodiments of the present application, the triple smelting process comprises: A: according to the alloy composition, smelting in a vacuum induction furnace, controlling the full melting temperature at 1510°C~1580°C, controlling the refining temperature at 1510-1580°C, and casting at 1420-1480°C to obtain the vacuum induction electrode bar; B: subjecting the vacuum induction electrode bar prepared in step A to protective atmosphere electroslag remelting to obtain an electroslag remelting ingot; filling argon gas throughout the step B process for protection, and the flow rate of the argon gas is 100-200 L / min; C: subjecting the electroslag remelting ingot obtained in step B to vacuum self-consuming remelting to obtain a vacuum self-consuming ingot; the vacuum self-consuming melting rate is 2.7-3.8 kg / min.

[0052] In specific embodiments of the present application, the multi-stage homogenization annealing includes: keeping the alloy ingot at 490-510°C for more than 4 h, then heating to 990-1010°C and keeping for more than 8 h, then heating to 1110-1140°C and keeping for more than 15 h, then heating to 1150-1180°C and keeping for more than 25 h, then heating to 1190-1220°C and keeping for more than 30 h, and then furnace cooling.

[0053] Fig.1 is non-equilibrium solidification segregation characteristics such as the segregation phase and element segregation of the alloy system (such as Embodiment 1) of the embodiment of the present application. Based on the characteristics of the alloy, the present application proposes a specific multi-stage homogenization annealing process, referring to the temperature curve of Fig.2, each stage of which corresponds to the elimination of the non-equilibrium segregation and / or element segregation of each temperature section to obtain a homogenized low-segregation alloy ingot. The annealing treatment at 1110-1140°C and 1150-1180°C helps to eliminate the low-melting point phase, and the annealing treatment at 1190-1220°C helps to eliminate the segregation of the ingot composition and improve the hot working performance of the alloy.

[0054] For example, in different embodiments, the multi-stage homogenization annealing may include: the alloy ingot are kept at 490°C, 495°C, 500°C, 505°C or 510°C for more than 4 h, then heating to 990°C, 995°C, 1000°C, 1005°C or 1010°C and keeping for more than 8 h, and then heating to 1110°C, 1115°C, 1120°C, 1125°C, 1130°C, 1135°C or 1140°C and keeping for more than 25 h, then heating to 1150°C, 1155°C, 1160°C, 1165C, 1170°C, 1175°C or 1180°C and keeping for more than 25 h, then heating to 1190°C, 1195°C, 1200°C, 1205°C, 1210°C, 1215°C or 1220°C and keeping for more than 30 h, and then cooling in the furnace.

[0055] In actual operation, the heating rate between adjacent insulation temperatures may be 2~120°C / h, such as 2°C / h, 10°C / h, 20°C / h, 40°C / h, 50°C / h, 60°C / h, 80°C / h, 100°C / h, 120°C / h or a range consisting of any two of them.

[0056] In specific embodiments of the present application, 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; the heating rate from 1150-1180°C to 1190-1220°C can be 3-7°C / h.

[0057] In specific embodiments of the present application, adopting a fast-forging machine to perform the upsetting and drawing cogging, and the deformation amount of each of the upsetting is 10%~55%, and the deformation amount of each of the drawing is 15%~60%. Furthermore, in the process of the upsetting and drawing cogging, the forging temperature is reduced one by one in the range of 1060 to 1180°C.

[0058] In actual operation, when adopting the fast-forging machine for upsetting and drawing cogging, heating the annealed ingot to 1060-1180°C, wrapping the ingot with insulation cotton, after the heat preservation is over, taking out of the furnace for forging, repeatedly upsetting and drawing the ingot by the fast-forging machine.

[0059] For example, in different embodiments, the deformation amount of each of the upsetting can be 10%, 20%, 30%, 40%, 50%, 55% or a range consisting of any two thereof; the deformation amount of each of the drawing can be 15%, 20%, 30%, 40%, 50%, 60% or a range consisting of any two thereof.

[0060] In specific embodiments of the present application, in the upsetting and drawing cogging, the time for heat preservation for returning to the furnace after each deformation is 1~6 h, which can achieve static recrystallization and refine grains, thereby improving thermoplasticity.

[0061] In specific embodiments of the present application, during the process of the upsetting and drawing cogging, performing forging in three forging temperature ranges with forging temperature sequentially reduced one by one; wherein, the first forging temperature is 1170-1180°C, the second forging temperature is 1150-1160°C, and the third forging temperature is 1130-1140°C.

[0062] In actual operation, the final forging temperature of the upsetting and drawing cogging is ≥950°C.

[0063] Sequentially reducing the forging temperature during the forging process of the present application, the forging temperature range is a two-phase zone. During the forging process, gradually coarsening the γ' phase to the micron level, the micron-level γ' phase reduces the deformation resistance of the alloy, promotes the dynamic recrystallization of the alloy, and then prepares fine-grained bars, thereby improving the thermoplasticity of the alloy.

[0064] In specific embodiments of the present application, during the process of the upsetting and drawing cogging, the upsetting deformation amount gradually increases, that is, the upsetting deformation amount at the first forging temperature is smaller than that at the second forging temperature, and the upsetting deformation amount at the second forging temperature is smaller than that at the third forging temperature.

[0065] In specific embodiments of the present application, in the radial forging treatment, a multi-pass forging in one heat is used to prepare an alloy of preset specifications. Furthermore, in the radial forging treatment, the keeping time at 1020°C -1150°C is 30-300 min.

[0066] In actual operation, before radial forging the intermediate billet after upsetting and drawing cogging, preheating the hammer head of the radial forging machine at 350-650°C, and the preheating time is ≥4h. During the radial forging treatment, returning the intermediate billet obtained by upsetting and drawing cogging to the furnace and wrapping, and keeping at 1020°C, 1040°C, 1050°C, 1060°C, 1080°C, 1100°C, 1120°C or 1150°C for 30min, 60min, 90min, 120min, 150min, 180min, 240min or 300min, and then forging once into a bar of preset specifications in the radial forging machine, and then air-cooling to room temperature after forging. The radial forging process ensures that the outer edge of the bar is fully deformed and achieves dynamic recrystallization to obtain an equiaxed crystal structure.

[0067] In specific embodiments of the present application, the method further includes: performing heat treatment of the alloy after the radial forging treatment, and the heat treatment includes solid solution treatment and aging treatment. Further, the solid solution treatment includes: treating by keeping at 1100-1150°C for 2-6 h; the aging treatment includes: treating by keeping at 800-900°C for 2-6 h and then air cooling and then treating by keeping at 750-820°C for 8-24 h and then air cooling.

[0068] During the solid solution treatment, fully re-smelting the γ' phase and cooperating with the subsequent primary aging treatment to form a γ' phase with a size of 100~500µm. After the secondary aging treatment, a uniformly dispersed γ' phase with a size of 10~50µm can be formed.

[0069] In specific embodiments of the present application, the average grain size of the heat-treated alloy is Grade 8 or above, and the grain size varation is less than 2 grades. Furthermore, the content of γ' phase of the heat-treated alloy is 50wt%~58wt%.

[0070] In the nickel-based wrought superalloy prepared by the method of the present application, the average grain size can reach above Grade 8, such as Grade 8, Grade 8.5, Grade 9, Grade 9.5 or any two of them. The grain size difference at each location of the bar is less than level 2, such as the grain size difference at the center, R / 2, and the edge is less than Grade 2, such as Grade 1.5, Grade 1, Grade 0.5, Grade 0 or any two of them.

[0071] The alloy of the present application has a high γ' phase content and sufficient strength. For example, in different embodiments, the γ' phase content can be 50wt%, 51wt%, 52wt%, 53wt%, 54wt%, 55wt%, 56wt%, 57wt%, 58wt% or a range consisting of any two thereof.

[0072] In specific embodiments of the present application, the tensile properties of the alloy after heat treatment satisfy: The tensile strength at room temperature is ≥1510MPa, the yield strength at room temperature is ≥ 1130MPa.

[0073] The tensile strength at 750°C is ≥1110MPa, the yield strength at 750°C is ≥960MPa.

[0074] The tensile strength at 800°C is ≥1000MPa, and the yield strength at 800°C is ≥900MPa.

[0075] In different embodiments, for example, the tensile properties of the alloy after heat treatment can be as follows: The tensile strength at room temperature may be 1510MPa, 1530MPa, 1550MPa, 1580MPa, 1600MPa, 1620MPa, 1650MPa, 1680MPa or a range consisting of any two thereof; the yield strength at room temperature may be 1130MPa, 1150MPa, 1180MPa, 1200MPa, 1220MPa, 1255MPa or a range consisting of any two thereof.

[0076] The tensile strength at 750°C may be 1110MPa, 1150MPa, 1160MPa, 1180MPa, 1200MPa, 1220MPa, 1240MPa, 1260MPa, 1270MPa, 1275MPa or a range consisting of any two of them; the yield strength at 750°C may be 960MPa, 980MPa, 1000MPa, 1020MPa, 1050MPa, 1080MPa, 1090MPa or a range consisting of any two of them.

[0077] The tensile strength at 800°C may be 1000MPa, 1020MPa, 1050MPa, 1080MPa, 1100MPa, 1120MPa, 1130MPa or a range consisting of any two thereof; the yield strength at 800°C may be 900MPa, 910MPa, 920MPa, 930MPa, 940MPa, 950MPa or a range consisting of any two thereof.

[0078] In specific embodiments of the present application, the creep rupture life of the heat-treated alloy satisfies: The creep rupture life of 60h minimum at 750°C under 620MPa.

[0079] The creep rupture life of 45h minimum at 800°C under 500MPa.

[0080] In different embodiments, the creep rupture life of the alloy after heat treatment can be exemplified as follows: The creep rupture life at 750°C / 620MPa may be 60h, 70h, 80h, 100h, 110h, 120h, 130h, 140h, 150h, 160h, 164h or a range consisting of any two thereof.

[0081] The creep rupture life at 800°C / 500MPa may be 45h, 50h, 60h, 70h, 80h, 85h or a range consisting of any two of them.

[0082] Another aspect of the present application provides a turbine disk, which is made of any one of the above-mentioned nickel-based wrought superalloys.Embodiment 1

[0083] The present embodiment provides a nickel-based wrought superalloy, which 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 0.06%, La 0.02%, Ce 0.02%, Nd 0.05%, Mg 0.03%, P 0.015%, S 0.0006%, and the rest is Ni.

[0084] A method for preparing the nickel-based wrought superalloy comprises the following steps: (1) Preparing a self-consuming ingot with a specification of Φ508 mm by adopting a triple smelting process of vacuum induction melting + protective atmosphere electroslag remelting + vacuum self-consuming remelting according to the alloy composition, polishing the surface of the self-consuming ingot, and spraying a conventional anti-oxidation coating.

[0085] Wherein, during the vacuum induction melting, adding the rare earth elements Sc, La, Ce and Nd in the form of intermediate alloys Ni-Sc, Ni-La, Ni-Ce and Ni-Nd at a low temperature refining stage of 1480~1510°C; controlling the full melting temperature to be 1520~1550°C, the refining temperature is 1520~1540°C, and casting at 1440~1450°C.

[0086] During the protective atmosphere electroslag remelting, the argon flow rate is 150±10L / min.

[0087] During the vacuum self-consuming remelting, the vacuum self-consuming melting rate is 3.2kg / min.

[0088] (2) Subjecting the consumable ingot obtained in step (1) to multi-stage homogenization annealing to obtain an annealed ingot with a length denoted by L; wherein, the multi-stage homogenization annealing comprises: heating to 500°C at 90°C / h and keeping for 4h, then heating to 1000°C at 80°C / h and keeping for 8h, then heating to 1140°C at 50°C / h and keeping for 15h, then heating to 1180°C at 10°C / h and keeping for 25h, then heating to 1210°C at 5°C / h and keeping for 30h, and then furnace cooling.

[0089] (3) Subjecting the annealed ingot obtained in step (2) to upsetting and drawing cogging and radial forging treatment, which specifically comprises the following steps: ① keeping the annealed ingot at 1180°C for 5h, taking out of the furnace and upsetting with a pressing amount of 20% and a pressing time of 40s to obtain a primary intermediate billet with a length of 0.8L; ② keeping the primary intermediate billet at 1180°C for 5h, taking out of the furnace and drawing for 90s to obtain a secondary intermediate billet with a length of L; ③ keeping the secondary intermediate billet at 1160°C for 4h, taking out of the furnace and upsetting with a pressing amount of 30% and a pressing time of 45s to obtain a tertiary intermediate billet with a length of 0.7L; ④ keeping the tertiary intermediate billet at 1160°C for 4h, taking it out of the furnace and drawing it for 100s to obtain a fourth intermediate billet with a length of 1.2L; ⑤ keeping the fourth intermediate billet at 1140°C for 4h, taking it out of the furnace and upsetting with a pressing amount of 40% and a pressing time of 50s to obtain a fifth intermediate billet with a length of 0.6L; ⑥ keeping the fifth intermediate billet at 1140°C for 4h, taking it out of the furnace and drawing it for 120s to obtain the sixth intermediate billet with a size of Φ235mm±10mm; ⑦ preheating the hammer head of the radial forging machine at 500°C for 4h; keeping the six intermediate billet at 1130°C for 5h, forging on the radial forging machine after taking out of the furnace, and forging to a size of Φ165mm±10mm, air cooling to room temperature, and machining to remove the black skin on the surface to obtain a Φ150mm finished bar.

[0090] (4) Heat treating the finished bar, and the heat treatment includes: keeping at 1150°C for 4 h and air cooling to room temperature; then keeping at 860°C for 6 h and air cooling; and then keeping at 780°C for 24 h and air cooling to room temperature.Embodiment 2

[0091] The present embodiment provides a nickel-based wrought superalloy, which comprises 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 the rest is Ni.

[0092] A method for preparing the nickel-based wrought superalloy comprises the following steps: (1) Preparing a self-consuming ingot with a specification of Φ508 mm by adopting a triple smelting process of vacuum induction melting + protective atmosphere electroslag remelting + vacuum self-consuming remelting according to the alloy composition, polishing the surface of the self-consuming ingot, and spraying a conventional anti-oxidation coating.

[0093] Wherein, during the vacuum induction melting, adding the rare earth elements Sc, La, Ce and Nd in the form of intermediate alloys Ni-Sc, Ni-La, Ni-Ce and Ni-Nd at a low temperature refining stage of 1480-1510°C; controlling the full melting temperature to be 1540~1550°C, the refining temperature is 1540-1550°C, and casting at 1440-1450°C.

[0094] During the protective atmosphere electroslag remelting, the argon flow rate is 150±10L / min.

[0095] During the vacuum self-consuming remelting, the vacuum self-consuming melting rate is 3.2kg / min.

[0096] (2) Subjecting the consumable ingot obtained in step (1) to multi-stage homogenization annealing to obtain an annealed ingot with a length denoted by L; wherein, the multi-stage homogenization annealing comprises: heating to 500°C at 90°C / h and keeping for 4h, then heating to 1000°C at 80°C / h and keeping for 8h, then heating to 1130°C at 50°C / h and keeping for 20h, then heating to 1170°C at 10°C / h and keeping for 30h, then heating to 1200°C at 5°C / h and keeping for 35h, and then furnace cooling.

[0097] (3) Subjecting the annealed ingot obtained in step (2) to upsetting and drawing cogging and radial forging treatment, which specifically comprises the following steps: ① keeping the annealed ingot at 1170°C for 5h, taking out of the furnace and upsetting with a pressing amount of 20% and a pressing time of 40s to obtain a primary intermediate billet with a length of 0.8L; ② keeping the primary intermediate billet at 1170°C for 5h, taking out of the furnace and drawing for 90s to obtain a secondary intermediate billet with a length of L; ③ keeping the secondary intermediate billet at 1150°C for 4h, taking out of the furnace and upsetting with a pressing amount of 30% and a pressing time of 45s to obtain a tertiary intermediate billet with a length of 0.7L; ④ keeping the tertiary intermediate billet at 1150°C for 4h, taking it out of the furnace and drawing it for 100s to obtain a fourth intermediate billet with a length of 1.2L; ⑤ keeping the fourth intermediate billet at 1130°C for 4h, taking it out of the furnace and upsetting with a pressing amount of 40% and a pressing time of 50s to obtain a fifth intermediate billet with a length of 0.6L; ⑥ keeping the fifth intermediate billet at 1130°C for 4h, taking it out of the furnace and drawing it for 120s to obtain the sixth intermediate billet with a size of Φ320mm±10mm; ⑦ preheating the hammer head of the radial forging machine at 500°C for 4h; keeping the six intermediate billet at 1120°C for 5h, forging on the radial forging machine after taking out of the furnace, and forging to a size of Φ265mm±10mm, air cooling to room temperature, and machining to remove the black skin on the surface to obtain a Φ250mm finished bar.

[0098] (4) Heat treating the finished bar, and the heat treatment includes: keeping at 1150°C for 4 h and air cooling to room temperature; then keeping at 860°C for 6 h and air cooling; and then keeping at 780°C for 24 h and air cooling to room temperature.Embodiment 3

[0099] The present embodiment provides a nickel-based wrought superalloy, which comprises 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 the rest is Ni.

[0100] A method for preparing the nickel-based wrought superalloy comprises the following steps: (1) Preparing a self-consuming ingot with a specification of Φ508 mm by adopting a triple smelting process of vacuum induction melting + protective atmosphere electroslag remelting + vacuum self-consuming remelting according to the alloy composition, polishing the surface of the self-consuming ingot, and spraying a conventional anti-oxidation coating.

[0101] Wherein, during the vacuum induction melting, adding the rare earth elements Sc, La, Ce and Nd in the form of intermediate alloys Ni-Sc, Ni-La, Ni-Ce and Ni-Nd at a low temperature refining stage of 1480-1510°C; controlling the full melting temperature to be 1540~1550°C, the refining temperature is 1540-1550°C, and casting at 1440-1450°C.

[0102] During the protective atmosphere electroslag remelting, the argon flow rate is 150±10L / min.

[0103] During the vacuum self-consuming remelting, the vacuum self-consuming melting rate is 3.2kg / min.

[0104] (2) Subjecting the consumable ingot obtained in step (1) to multi-stage homogenization annealing to obtain an annealed ingot with a length denoted by L; wherein, the multi-stage homogenization annealing comprises: heating to 500°C at 90°C / h and keeping for 4h, then heating to 1000°C at 80°C / h and keeping for 8h, then heating to 1120°C at 50°C / h and keeping for 25h, then heating to 1160°C at 10°C / h and keeping for 25h, then heating to 1190°C at 5°C / h and keeping for 38h, and then furnace cooling.

[0105] (3) Subjecting the annealed ingot obtained in step (2) to upsetting and drawing cogging and radial forging treatment, which specifically comprises the following steps: ① keeping the annealed ingot at 1180°C for 5h, taking out of the furnace and upsetting with a pressing amount of 20% and a pressing time of 40s to obtain a primary intermediate billet with a length of 0.8L; ② keeping the primary intermediate billet at 1180°C for 4h, taking out of the furnace and drawing for 90s to obtain a secondary intermediate billet with a length of L; ③ keeping the secondary intermediate billet at 1160°C for 4h, taking out of the furnace and upsetting with a pressing amount of 30% and a pressing time of 45s to obtain a tertiary intermediate billet with a length of 0.7L; ④ keeping the tertiary intermediate billet at 1160°C for 4h, taking it out of the furnace and drawing it for 100s to obtain a fourth intermediate billet with a length of 1.2L; ⑤ keeping the fourth intermediate billet at 1140 °C for 4h, taking it out of the furnace and upsetting with a pressing amount of 40% and a pressing time of 50s to obtain a fifth intermediate billet with a length of 0.6L; ⑥ keeping the fifth intermediate billet at 1140°C for 4h, taking it out of the furnace and drawing it for 120s to obtain the sixth intermediate billet with a size of Φ320mm±10mm; ⑦keeping the sixth intermediate billet at 1130°C for 4h, taking it out of the furnace and drawing it for 90s to obtain the seventh intermediate billet with a size of Φ280mm±10mm; ⑧ preheating the hammer head of the radial forging machine at 500°C for 4h; keeping the seventh intermediate billet at 1100°C for 5h, forging on the radial forging machine after taking out of the furnace, and forging to a size of Φ230mm±10mm, air cooling to room temperature, and machining to remove the black skin on the surface to obtain a Φ200mm finished bar.

[0106] (4) Heat treating the finished bar, and the heat treatment includes: keeping at 1150°C for 4 h and air cooling to room temperature; then keeping at 860°C for 6 h and air cooling; and then keeping at 780°C for 24 h and air cooling to room temperatureEmbodiment 4

[0107] This embodiment refers to the nickel-based wrought superalloy and its preparation method of Embodiment 1, the only difference is that the alloy composition is different.

[0108] The nickel-based wrought superalloy of the present embodiment includes 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 the rest is Ni balance.Embodiment 5

[0109] This embodiment refers to the nickel-based wrought superalloy and its preparation method of Embodiment 1, the only difference is that the alloy composition is different.

[0110] The nickel-based wrought superalloy of the present embodiment includes 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.Embodiment 6

[0111] This embodiment refers to the nickel-based wrought superalloy and its preparation method of Embodiment 1, the only difference is that the alloy composition is different.

[0112] The nickel-based wrought superalloy of the present embodiment includes 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 the rest is Ni.Embodiment 7

[0113] This embodiment refers to the nickel-based wrought superalloy and its preparation method of Embodiment 1, the only difference is that the homogenization annealing process in step (2) in the preparation method is different.

[0114] The homogenization annealing of the present embodiment comprises: heating to 500°C at 90°C / h and keeping for 4h, then heating to 1000°C at 80°C / h and keeping for 8h, then heating to 1150°C at 50°C / h and keeping for 15h, then heating to 1200°C at 5°C / h and keeping for 25h, and then furnace cooling.Embodiment 8

[0115] This embodiment refers to the nickel-based wrought superalloy and its preparation method of Embodiment 1, the only difference is that step (3) in the preparation method is different.

[0116] The step (3) of the present embodiment includes: subjecting the annealed ingot obtained in step (2) to upsetting and drawing cogging and radial forging treatment, which specifically comprises the following steps: ① keeping the annealed ingot at 1180°C for 5h, taking out of the furnace and upsetting with a pressing amount of 20% and a pressing time of 40s to obtain a primary intermediate billet with a length of 0.8L; ② keeping the primary intermediate billet at 1180°C for 5h, taking out of the furnace and drawing for 90s to obtain a secondary intermediate billet with a length of L; ③ keeping the secondary intermediate billet at 1180°C for 4h, taking out of the furnace and upsetting with a pressing amount of 30% and a pressing time of 45s to obtain a tertiary intermediate billet with a length of 0.7L; ④ keeping the tertiary intermediate billet at 1180°C for 4h, taking it out of the furnace and drawing it for 100s to obtain a fourth intermediate billet with a length of 1.2L; ⑤ keeping the fourth intermediate billet at 1180°C for 4h, taking it out of the furnace and upsetting with a pressing amount of 40% and a pressing time of 50s to obtain a fifth intermediate billet with a length of 0.6L; ⑥ keeping the fifth intermediate billet at 1180°C for 4h, taking it out of the furnace and drawing it for 120s to obtain the sixth intermediate billet with a size of Φ235mm±10mm; ⑦ preheating the hammer head of the radial forging machine at 500°C for 4h; keeping the sixth intermediate billet at 1130°C for 5h, forging on the radial forging machine after taking out of the furnace, and forging to a size of Φ165mm±10mm, air cooling to room temperature, and machining to remove the black skin on the surface to obtain a Φ150mm finished bar. Comparative embodiment 1

[0117] Comparative embodiment 1 refers to the nickel-based wrought superalloy and its preparation method of Embodiment 1, the difference is that the alloy composition is different.

[0118] The nickel-based wrought superalloy of Comparative embodiment 1 includes 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 the rest is Ni.Comparative embodiment 2

[0119] Comparative embodiment 2 refers to the nickel-based wrought superalloy and its preparation method of Embodiment 1, the difference is that the alloy composition is different.

[0120] The nickel-based wrought superalloy of Comparative embodiment 2 includes 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 the rest is Ni.Comparative embodiment 3

[0121] Comparative embodiment 3 refers to the nickel-based wrought superalloy and its preparation method of Embodiment 1, the difference is that the homogenization annealing process in step (2) in the preparation method is different.

[0122] The homogenization annealing of Comparative embodiment 3 comprises: heating to 1210°C at 5°C / h and keeping for 30h, and then furnace cooling.Comparative embodiment 4

[0123] Comparative embodiment 4 refers to the nickel-based wrought superalloy and its preparation method of Embodiment 2, the difference is that step (3) in the preparation method does not include radial forging treatment.

[0124] Step (3) of Comparative embodiment 4 comprises: preparing a sixth intermediate billet according to ①~⑥ of step (3) of Embodiment 2; then keeping the sixth intermediate billet at 1120°C for 5 h, drawing for 120 s after taking out of the furnace, drawing and rounding to Φ280 mm±10 mm, air cooling to room temperature, machining and polishing to remove the black skin on the surface to obtain a finished bar with a specification of Φ250 mm.Experimental embodiment

[0125] Respectively subjecting the bar sample prepared in each embodiment and comparative embodiment to metallographic observation and mechanical property test. The grain structures are shown in Table 1, and the mechanical properties are shown in Table 2. Fig.3 and Fig.4 are grain structure diagrams of the bars prepared in Embodiment 1 and Comparative embodiment 1 of the present application, respectively. Figure 5 is a microstructure diagram of the annealed alloy ingot prepared in Comparative embodiment 1; Fig. 6 and Fig.7 are a tensile fracture morphology at 800°C and a fracture morphology of long-term specimen under an 800°C. / 500 MPa environment of the bar prepared in Embodiment 1 of the present application. Figures 8 and 9 are a tensile fracture morphology at 800°C and a fracture morphology of long-term specimen under an 800°C. / 500 MPa environment of the bar prepared in Comparative embodiment 1 of the present application. Table 1 Metallographic structures of different barsNumberCenter grain sizeR / 2 grain sizeEdge grain sizeStructure homogeneityThe content of y' phase (wt%)Embodiment 1Grade 9Grade 9Grade 9.5Homogeneous and fine grain57Embodiment 2Grade 8.5Grade 8.5Grade 9Homogeneous and fine grain56Embodiment 3Grade 9Grade 9Grade 9.5Homogeneous and fine grain56Embodiment 4Grade 6.5Grade 8Grade 8.5Fine grain, local mixed grain56Embodiment 5Grade 7Grade 9Grade 8.5Fine grain, local mixed grain55Embodiment 6Grade 6.5Grade 8.5Grade 8.5Fine grain, local mixed grain55Embodiment 7Grade 6Grade 7.5Grade 8Fine grain, local mixed grain57Embodiment 8Grade 5.5Grade 6Grade 7Mixed grain57Comparative embodiment 1Grade 5Grade 6.5Grade 8Mixed grain55Comparative embodiment 2Grade 4.5Grade 6.5Grade 7Mixed grain55Comparative embodiment 3Grade 4.5Grade 7Grade 8Mixed grain55Comparative embodiment 4Grade 4Grade 5.5Grade 5Mixed grain55 Table 2 Mechanical properties test results of different bars NumberMechanical properties at room temperature / MPaMechanical properties at 750 °C / MPaMechanical properties at 800°C / MPaCreep rupture life at 750°C / 620MPaCreep rupture life at 800°C / 500MPaTensile strengthYield strengthTensile strengthYield strengthTensile strengthYield strengthEmbodiment 11680125512751090113095014578Embodiment 21655124712601076112093916485Embodiment 31675125312681080113294815272Embodiment 415721172116599510909157854Embodiment 515831156113898810829026951Embodiment 615661145115499210659077148Embodiment 715231140113298010729036347Embodiment 815101130113596810589006850Comparative embodiment 11385108810748609528253914Comparative embodiment 21415110511028859708405628Comparative embodiment 31350106510688359328053925Comparative embodiment 41364107010858509408153126

[0126] It can be seen from the test results that the nickel-based wrought superalloy of the present application has the characteristics of high homogeneity and low segregation, the average grain size of the fine-grained bar is finer than Grade 8, and the grain size variation is less than 2 grades; the tensile strength at room temperature is ≥ 1655MPa, the tensile strength at 750°C is ≥ 1260MPa, the creep rupture life at 750°C / 620MPa is ≥145h, and the creep rupture life at 800°C / 500MPa is ≥72h.

[0127] Finally, it should be noted that the above embodiments are only used to illustrate the technical solutions of the present application rather than to limit it. Although the present application has been described in detail with reference to the above-mentioned embodiments, those skilled in the art should understand that they can still modify the technical solutions described in the above-mentioned embodiments, or perform the equivalent substitution of some or all of the technical features. However, these modifications or replacements do not cause the essence of the corresponding technical solutions to deviate from the scope of the technical solutions of the embodiments of the present application.

Claims

1. A high-strength high-toughness nickel-based wrought superalloy resistant to 800°C, characterized in that, the nickel-based wrought superalloy comprises 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%, and the rest is Ni and unavoidable impurities.

2. The high-strength and high-toughness nickel-based wrought superalloy resistant to 800°C according to claim 1, characterized in that, a sum of mass percentages of Al, Ti and Nb satisfies: Al+Ti+Nb≥9.5%; and a mass ratio Al / Ti of Al to Ti is 1.15-1.45.

3. The high-strength and high-toughness nickel-based wrought superalloy resistant to 800°C according to claim 1, characterized in that, a sum of mass percentages of Mo and W satisfies: Mo+W≥7.5%; and a mass ratio Mo / W of Mo to W is 1.3-1.45.

4. The high-strength and high-toughness nickel-based wrought superalloys resistant to 800 °C according to claim 1, characterized in that, a sum of mass percentages of C, B and Zr satisfies: C+B+Zr≤0.15%, and / or, a sum of mass percentages of Sc, La, Ce and Nd satisfies: Sc+La+Ce+Nd≤0.2%.

5. A preparation method of the high-strength and high-toughness nickel-based wrought superalloys resistant to 800 °C according to any one of claims 1~4, characterized in that, the preparation method comprises the following steps: (a) preparing an alloy ingot by a triple smelting process according to alloy components; (b) subjecting the alloy ingot to a multi-stage homogenization at 1100-1220°C to obtain an annealed ingot; (c) subjecting the annealed ingot to an upsetting and drawing at 1060-1180°C, and then subjecting to a radial forging at 1020-1150°C.

6. The preparation method of the high-strength and high-toughness nickel-based wrought superalloys resistant to 800°C according to claim 5, characterized in that, the multi-stage homogenization comprises: keeping the alloy ingot at 490-510°C for more than 4 h, then heating to 990-1010°C and keeping for more than 8 h, then heating to 1110-1140°C and keeping for more than 15 h, then heating to 1150-1180°C and keeping for more than 25 h, then heating to 1190-1220°C and keeping for more than 30 h, and then furnace cooling.

7. The preparation method of the high-strength and high-toughness nickel-based wrought superalloys resistant to 800°C according to claim 5, characterized in that, adopting a high-speed forging press to perform the upsetting and drawing cogging, and a deformation amount of each of the upsetting is 10%~55%, and a deformation amount of each of the drawing is 15%~60%; preferably, a sequential heat temperature decrease from 1060~1180°C was implemented in the process of the upsetting and drawing; preferably, in the radial forging, adopting a multi-pass forging in one heat to prepare an alloy of a preset specification; preferably, in the radial forging, a holding time at 1020°C-1150°C is 30-300 min.

8. The preparation method of the high-strength and high-toughness nickel-based wrought superalloys resistant to 800°C according to claim 5, characterized in that, the preparation method also comprises: performing heat treatment of an alloy after the radial forging to obtain a heat-treated alloy, and the heat treatment comprises a solid solution treatment and an aging; preferably, the solid solution treatment comprises: holding at 1100-1150°C for 2-6 h; the aging comprises: holding at 800-900°C for 2-6 h and air cooled, and then holding at 750-820°C for 8-24 h and then air cooled.

9. The preparation method of the high-strength and high-toughness nickel-based wrought superalloys resistant to 800°C according to claim 8, characterized in that, the heat-treated alloy has at least one of characteristics as follow: (1) an average grain size of the heat-treated alloy is above Grade 8, and a grain size varation is less than 2 grades; (2) a content of a γ' phase is 50wt%~58wt%; (3) a tensile strength at room temperature is ≥1510MPa, a yield strength at room temperature is ≥1130MPa; (4) a tensile strength at 750°C is ≥1110MPa, a yield strength at 750°C is ≥960MPa; (5) a tensile strength at 800°C is ≥1000MPa, a yield strength at 800°C is ≥900MPa; (6) the creep rupture life of 60h minimum at 750°C under 620MPa; (7) the creep rupture life of 45h minimum at 800°C under 500MPa.

10. A turbine disk, characterized in that, the turbine disk is made of the high-strength and high-toughness nickel-based wrought superalloys resistant to 800 °C mentioned in any one of claims 1~4 or the high-strength and high-toughness nickel-based wrought superalloys resistant to 800 °C obtained by the preparation method mentioned in any one of claims 5~9.

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