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 superalloys, and particularly relates to an 800°C-resistant high-strength and high-toughness nickel-based wrought superalloy and a preparation method therefor, and a turbine disc. The nickel-based wrought superalloy comprises: 10-12% of Cr, 14-16% of Co, 2.5-3.5% of W, 3.75-5% of Mo, 3.5-4.5% of Al, 2.5-3.1% of Ti, 3-3.8% of Nb, 0.01-0.08% of C, 0.005-0.05% of B, 0.01-1% of V, 0.025-0.075% of Zr, 0.005-0.3% of Si, 0.02-0.05% of Re, 0.001-0.08% of Sc, 0.001-0.08% of La, 0.001-0.05% of Ce, 0.001-0.05% of Nd, 0.001-0.15% of Mg, and the balance of Ni. In the present application, by rationally designing the components of the nickel-based superalloy in cooperation, and by coarsening a γ' phase by means of process regulation and control, the alloy achieves good hot working performance for engineering, and a turbine disc satisfying requirements can also be obtained by means of a heat treatment.

Description

800℃ resistant high-strength and tough nickel-based deformable high-temperature alloy, preparation method thereof, and turbine disk Technical Field

[0001] The present application relates to the technical field of high-temperature alloys, and in particular to a high-strength and tough nickel-based deformable high-temperature alloy resistant to 800°C, a preparation method thereof, and a turbine disk. Background Art

[0002] Turbine disks are important hot-end components of aircraft engines. Operating in harsh environments, they are primarily made of deformed high-temperature alloys. With the development of the aviation industry, the outlet temperature of aircraft engines has gradually increased, so the temperature-bearing capacity of nickel-based deformed high-temperature alloys used in turbine disks has also increased. These alloys are required to have excellent medium-temperature (600-800°C) service performance. The main design methods for strengthening and toughening deformed high-temperature alloys include: precipitation strengthening by adding Al, Ti, and Nb elements to form γ" or γ' phases; solid solution strengthening by adding W and Mo elements; and grain boundary strengthening and toughening by adding microalloying elements such as B, Zr, and C.

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

[0004] At present, the main typical grades of nickel-based deformable high-temperature alloys are GH4169 alloy resistant to 650℃, GH169D alloy resistant to 700℃, GH4720Li alloy resistant to 750℃ and GH4065A alloy. The alloying degree of these alloys gradually increases. The high alloying degree will significantly reduce the plasticity of the alloy in processing (high temperature) and service (medium temperature) environments; the higher the alloying degree, the more serious the element segregation of the high-temperature alloy ingots and the more complex the solidification segregation phase. The element segregation of large-sized alloy ingots is serious, and the segregation structure inheritance is difficult to eliminate. The ingot specifications are limited to small ingots, the hot working window is narrow, the deformation resistance is large, and it is very easy to crack. These problems make the process and structure control of the blanking process extremely difficult, the yield rate is extremely low, and the structure is uneven, which has become a key problem restricting the preparation and application of high-quality bars of high-alloyed high-temperature alloys.

[0005] In view of this, this application is hereby filed. Summary of the Invention

[0006] One purpose of the present application is to provide a high-strength and tough nickel-based deformable high-temperature alloy that is resistant to 800°C. It not only has both strength and toughness in the medium temperature range, but also has high-temperature machinability, and can provide the necessary high-thermal-strength nickel-based deformable high-temperature material for high thrust ratio engines.

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

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

[0009] In order to achieve the above-mentioned object of the present application, the present application provides, on one hand, a nickel-based deformable high-temperature alloy resistant to 800°C with high strength and toughness, comprising the following components in percentage by mass:

[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% to 0.15%, the balance being Ni and unavoidable impurities.

[0011] In a specific embodiment 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 of Al to Ti, Al / Ti, is 1.15-1.45.

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

[0013] In a specific embodiment of the present application, 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 application, 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 application, the unavoidable impurities include any one or more of S, P, N, O, Cu, Ag, Si, Mn, Ca, Sn, and Pb. Furthermore, in the alloy, the impurity elements include, by mass percentage: 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] 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 deformable high-temperature alloys resistant to 800°C, comprising the following steps:

[0017] (a) preparing alloy ingots by a triple smelting process according to the alloy composition;

[0018] (b) subjecting the alloy ingot to multi-stage homogenization annealing at 1100-1220° C. to obtain an annealed ingot;

[0019] (c) The annealed ingot is subjected to upsetting and drawing at 1060-1180° C. to form a blank, and then subjected to radial forging at 1020-1150° C.

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

[0021] In a specific embodiment of the present application, a high-speed forging machine is used for the upsetting and drawing process, with the upsetting deformation of each fire being 10% to 55%, and the drawing deformation of each fire being 15% to 60%. Furthermore, during the upsetting and drawing process, the forging temperature is lowered in each fire within the range of 1060 to 1180°C.

[0022] In a specific embodiment of the present application, in the upsetting and drawing process, the time for returning to the furnace and keeping warm after each deformation is 1 to 6 hours.

[0023] In a specific embodiment of the present application, the radial forging process uses a one-fire-multiple-pass deformation method to prepare an alloy of preset specifications. Furthermore, the radial forging process is performed at a holding time of 1020-1150° C. for 30-300 minutes.

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

[0025] In a specific embodiment of the present application, the average grain size of the heat-treated alloy is above level 8, with a grain difference of less than level 2. Furthermore, the γ' phase content of the heat-treated alloy is 50 wt% to 58 wt%.

[0026] In a specific embodiment of the present application, the tensile properties of the heat-treated alloy meet the following requirements:

[0027] Room temperature tensile strength ≥1510MPa, room temperature yield strength ≥1130MPa;

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

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

[0030] In a specific embodiment of the present application, the durability of the alloy after heat treatment satisfies:

[0031] The endurance life of 750℃ / 620MPa is ≥60h;

[0032] The endurance life at 800℃ / 500MPa is ≥45h.

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

[0034] Compared with the prior art, the present invention has the following advantages:

[0035] (1) This application rationally designs the composition of nickel-based high-temperature alloys, adds certain rare earth elements and trace elements such as C, B, and Zr, improves the grain boundary bonding strength of the alloy, reduces or eliminates the brittle tendency 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 intermediate temperature brittleness and low plasticity of the alloy in the service temperature range, so that the alloy can have both high strength and toughness and good process performance;

[0036] (2) The present application adopts a multi-stage homogenization annealing process, a combined upsetting and radial forging process to produce a highly homogeneous and low-segregation alloy in accordance with the alloy composition, and ensures its high strength and toughness at 800°C. By controlling the process, the strengthening effect of the γ' phase is weakened, and the alloy is controlled to have a fine-grained structure, so that the alloy can be prepared by hot working in engineering.

[0037] (3) The nickel-based deformed high-temperature alloy of the present application meets the high thermal strength requirements of materials for high thrust ratio engines, which is of great significance to the development of the aviation field. BRIEF DESCRIPTION OF THE DRAWINGS

[0038] In order to more clearly illustrate the specific implementation methods of the present application or the technical solutions in the prior art, the following is a brief introduction to the drawings required for use in the specific implementation methods or the description of the prior art. Obviously, the drawings described below are some implementation methods of the present application. For ordinary technicians in this field, other drawings can be obtained based on these drawings without any creative work.

[0039] FIG1 shows the non-equilibrium solidification segregation characteristics such as alloy segregation phase and element segregation of the alloy system of the embodiment of the present application;

[0040] FIG2 is a multi-stage high-temperature diffusion homogenization annealing temperature curve provided in an embodiment of the present application;

[0041] FIG3 is a grain structure diagram of the rod prepared in Example 1 of the present application;

[0042] FIG4 is a grain structure diagram of the rod prepared in Comparative Example 1;

[0043] FIG5 is a microstructure diagram of the annealed alloy ingot prepared in Comparative Example 1;

[0044] FIG6 is a tensile fracture morphology of the rod prepared in Example 1 of the present application at 800° C.;

[0045] FIG7 is a fracture morphology of a long-lasting specimen of the rod prepared in Example 1 of the present application under an 800° C. / 500 MPa environment;

[0046] FIG8 is a tensile fracture morphology of the rod prepared in Comparative Example 1 at 800° C.;

[0047] FIG9 is a fracture morphology of a long-lasting specimen of the rod prepared in Comparative Example 1 under 800° C. / 500 MPa conditions. DETAILED DESCRIPTION

[0048] 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 the present application's embodiments, rather than all of the embodiments, and are only used to illustrate the present application, and should not be considered as limiting the scope of the present application. Based on the embodiments in the present application, all other embodiments obtained by those of ordinary skill in the art without making creative work are within the scope of protection of the present application. Those who do not specify specific conditions in the embodiments are carried out according to conventional conditions or the conditions recommended by the manufacturer. Those whose reagents or instruments are not specified as manufacturers are conventional products that can be purchased commercially.

[0049] On one hand, the present application provides a nickel-based deformable high-temperature alloy with high strength and toughness resistant to 800°C, comprising the following components in percentage by mass:

[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% to 0.15%, the balance being Ni and unavoidable impurities.

[0051] This application improves the hot working performance and mechanical properties of the alloy by rationally designing the composition of the nickel-based high-temperature alloy. The alloy of this application is a highly alloyed, difficult-to-deform nickel-based high-temperature alloy. The γ' phase content in the alloy is 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 very high, and excellent high-temperature strength can be obtained at 800°C. Rare earth elements such as Sc and grain boundary strengthening elements such as Zr are added to the alloy to improve the grain boundary bonding strength, and the isocratic transformation temperature between the grain boundary and the interior of the grain is shifted to a higher temperature, so that the alloy obtains better toughness and avoids medium-temperature brittleness at around 800°C. In conjunction with the process control during the hot working process, the strengthening effect of the γ' phase is weakened, and the alloy is controlled into a fine-grained structure, so that the alloy can be prepared by hot working in engineering.

[0052] A large amount of solid solution strengthening elements, precipitation strengthening elements, grain boundary strengthening elements and rare earth elements are added to the alloy of the present application, and the specific 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% to 58wt%, significantly higher than other high-heat-strength nickel-based deformable high-temperature alloys. The large amount of precipitation-strengthening phase ensures sufficient strength for the alloy. The Nb element enters the γ' phase to form Ni3 (Al, Ti, Nb), which increases the volume fraction of the γ' phase and the reverse domain boundary energy in the alloy, thereby improving the alloy's strength. To balance the alloy's strength and workability, the present application further preferably regulates 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 to 1.45.

[0054] 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 thereof; 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 thereof; 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 a range consisting of any two thereof; the mass ratio of Al and Ti, Al / 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.

[0055] Adding solid-solution strengthening elements like W and Mo to the alloy significantly distorts the lattice and reduces the stacking fault energy, thereby improving the alloy's yield strength and creep properties. However, excessive or inappropriately proportioned additions of W and Mo can lead to the formation of a harmful brittle TCP phase over extended use. Under thermal-mechanical coupling, cracks initiate and propagate at the TCP-matrix interface, deteriorating the alloy's performance. Therefore, the present invention further regulates the sum of the mass percentages of Mo and W to satisfy the following: W + Mo ≥ 7.5%, and the mass ratio of M to W, Mo / W, is between 1.3 and 1.45, achieving sufficient solid-solution strengthening.

[0056] 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%, or 5%. Furthermore, the sum of the mass percentages of Mo and W is 7.5% to 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 of Mo to W, Mo / 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.

[0057] The addition of C to the alloy of the present application causes the element to segregate to grain boundaries or form grain boundary carbides, which precipitate in the form of blocks or particles at the grain boundaries. This can inhibit abnormal grain growth, prevent grain boundary sliding, strengthen the grain boundaries, and improve long-term strength and plasticity. However, excessive addition causes the carbides to precipitate continuously along the grain boundaries in the form of thin films, embrittle the grain boundaries, and cause the alloy to exhibit notch sensitivity. Therefore, the C content in the present application is controlled within a range of 0.01% to 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 thereof.

[0058] Zr and B elements are added to the alloy of the present application, and the elements are segregated at the grain boundaries, thereby increasing the grain boundary bonding strength, reducing the interfacial energy, improving the grain boundary morphology, and improving the creep endurance strength and low-cycle fatigue performance. At the same time, the Zr element can also be used as a purifier to react with carbon sulfides or carbides of elements such as C and S in the alloy, thereby removing harmful impurity elements such as S and P; while adding excessive B element will cause boride precipitation at the grain boundaries or induce the generation of hot cracks. Therefore, in the present application, the Zr content is controlled at 0.025% to 0.075%, and the B content is controlled at 0.005% to 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.

[0059] The present application further preferably controls 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, endurance strength, and ductility while reducing or eliminating brittle tendencies in the service temperature range, avoiding grain boundary embrittlement and the generation of thermal cracks. 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 thereof.

[0060] The addition of Sc, La, Ce and Nd rare earth elements to the alloy of the present application can reduce the critical supercooling of dendrite nucleation, produce compositional supercooling at the solid-liquid interface front during the solidification process of the alloy, reduce the critical nucleation radius and nucleation work, increase the nucleation rate, thereby refining the dendrites and effectively reducing segregation. At the same time, rare earth elements can also remove O and purify grain boundaries, reduce or avoid the generation of intergranular cracks in the alloy in the medium temperature range. At the same time, 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 eliminate in the subsequent homogenization and thermal deformation process. Therefore, the rare earth content in the present application should be strictly controlled. In particular, the Sc content is controlled within 0.001% to 0.08%, the La content is controlled within 0.001% to 0.08%, the Ce content is controlled within 0.001% to 0.05%, and the Nd content is controlled in the range of 0.001% to 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.

[0061] The sum of the mass percentages of Sc, La, Ce, and Nd is further adjusted to satisfy the following: Sc+La+Ce+Nd≤0.2%, such as 0.1% to 0.2%, to optimize the aforementioned 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 thereof.

[0062] The Cr element in the alloy of the present application is the main forming element of carbides, and can also improve the oxidation resistance and corrosion resistance of the alloy. Too low a Cr content will lead to a decrease in the oxidation resistance and corrosion resistance of the alloy, and too low a Cr content will lead to the formation of 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% to 12%, and combines it with the composition of other elements so that the alloy system of the present application takes into account the improvement of oxidation resistance and corrosion resistance as well as plasticity and toughness. 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 composed of any two of them.

[0063] In a specific embodiment of the present application, unavoidable impurities include any one or more of S, P, N, O, Cu, Ag, Si, Mn, Ca, Sn, and Pb. Furthermore, in the alloy, the impurity elements include, by mass percentage: 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 application, the 800°C resistant high-strength and high-toughness nickel-based deformable high-temperature alloy includes 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%, 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% to 0.04%, the balance being Ni and unavoidable impurities.

[0066] 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 deformable high-temperature alloys resistant to 800°C, comprising the following steps:

[0067] (a) preparing alloy ingots by a triple smelting process according to the alloy composition;

[0068] (b) performing multi-stage homogenization annealing on the alloy ingot at 1100-1220° C. to obtain an annealed ingot;

[0069] (c) The annealed ingot is subjected to upsetting and drawing at 1060-1180°C to form a blank, and then subjected to radial forging at 1020-1150°C.

[0070] The material of the present application has a high degree of alloying and requires high purity of the alloy. Through the triple smelting process of vacuum induction melting + electroslag remelting + vacuum induction remelting, combined with the alloy composition of the present application, an ingot with higher purity can be obtained.

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

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

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

[0074] A: According to the alloy composition, the materials are 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 cast at 1420-1480℃.

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

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

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

[0078] Figure 1 shows the non-equilibrium solidification segregation characteristics of the alloy system (e.g., Example 1), including segregation phases and elemental segregation, for the alloy system of the present application. Based on these characteristics, the present application proposes a specific multi-stage homogenization annealing process. Referring to the temperature curve in Figure 2, each stage eliminates the non-equilibrium segregation and / or elemental segregation in each temperature range to obtain a homogenized low-segregation alloy ingot. Annealing treatments at 1110-1140°C and 1150-1180°C help eliminate low-melting-point phases, while annealing treatments at 1190-1220°C help eliminate ingot composition segregation and improve the alloy's hot workability.

[0079] For example, in different embodiments, the multi-stage homogenization annealing may include: subjecting the alloy ingot to a holding temperature of 490°C, 495°C, 500°C, 505°C or 510°C for more than 4 hours, then heating the alloy ingot to 990°C, 995°C, 1000°C, 1005°C or 1010°C and holding the temperature for more than 8 hours, and then heating the alloy ingot to 1110°C, 1115°C, 1120°C, 1125°C, 1130°C, 1140°C, 1155°C, 1160°C, 1170°C, 1180°C, 1190°C, 2200°C, 2210°C, 2225°C, 2230°C, 2240°C, 2255°C, 2260°C, 2270°C, 2280°C, 2290°C, 2310°C, 2320°C, 2330°C, 2340°C, 2360°C, 2370°C, 2380°C, 2390°C, 2400°C, 2410°C, 2420°C, 2430°C, 2440°C, 2450°C, 2460°C, 2470°C, 2480°C, 2490°C, Heat the temperature to 135℃ or 1140℃ and keep warm for more than 15 hours, then heat it to 1150℃, 1155℃, 1160℃, 1165℃, 1170℃, 1175℃ or 1180℃ and keep warm for more than 25 hours, then heat it to 1190℃, 1195℃, 1200℃, 1205℃, 1210℃, 1215℃ or 1220℃ and keep warm for more than 30 hours, then cool in the furnace.

[0080] In actual operation, the heating rate between adjacent holding temperatures may be 2 to 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 thereof.

[0081] In a specific embodiment 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.

[0082] In a specific embodiment of the present application, a high-speed forging machine is used for upsetting and drawing, with the upsetting deformation of each fire being 10% to 55%, and the elongation deformation of each fire being 15% to 60%. Furthermore, during the upsetting and drawing process, the forging temperature is lowered in the range of 1060 to 1180°C with each fire.

[0083] In actual operation, when using a fast forging machine for upsetting and drawing, the annealed ingot is heated to 1060-1180°C, and the ingot is wrapped with insulation cotton. After the insulation is completed, it is taken out of the furnace for forging, and the ingot is repeatedly upset and drawn using a fast forging machine.

[0084] For example, in different embodiments, the deformation amount of each upsetting fire can be 10%, 20%, 30%, 40%, 50%, 55% or a range consisting of any two thereof; the deformation amount of each drawing fire can be 15%, 20%, 30%, 40%, 50%, 60% or a range consisting of any two thereof.

[0085] In a specific embodiment of the present application, in the upsetting and drawing process, the time of returning to the furnace and holding the heat after each heat deformation is 1 to 6 hours. Returning to the furnace and holding the heat after each heat deformation can achieve static recrystallization, refine grains, and improve thermoplasticity.

[0086] In a specific embodiment of the present application, during the upsetting and drawing of the blank, forging can be performed within three forging temperature ranges with the forging temperature being lowered step by step; 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.

[0087] In actual operation, the final forging temperature of the upsetting and drawing billet is ≥950℃.

[0088] In the forging process of the present application in which the forging temperature is lowered in successive fires, the forging temperature range is a two-phase region. During the forging process, the γ' phase gradually coarsens 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 produces fine-grained rods, thereby improving the thermoplasticity of the alloy.

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

[0090] In a specific embodiment of the present application, during the radial forging process, a single-fire, multi-pass deformation method is used to prepare an alloy of predetermined specifications. Furthermore, during the radial forging process, the holding time at 1020-1150° C. is 30-300 minutes.

[0091] In practice, the hammer head of the radial forging machine is preheated to a temperature of 350-650°C for 4 hours or longer before radial forging the intermediate billet after upsetting and drawing. During radial forging, the intermediate billet obtained from upsetting and drawing is returned to the furnace and jacketed. It is then held at 1020°C, 1040°C, 1050°C, 1060°C, 1080°C, 1100°C, 1120°C, or 1150°C for 30 minutes, 60 minutes, 90 minutes, 120 minutes, 150 minutes, 180 minutes, 240 minutes, or 300 minutes. The bar is then forged in a radial forging machine to a predetermined specification and air-cooled to room temperature. This radial forging process ensures sufficient deformation of the outer edge of the bar, enabling dynamic recrystallization and resulting in an equiaxed grain structure.

[0092] In a specific embodiment of the present application, the alloy after the radial forging treatment is further subjected to heat treatment, which includes solution treatment and aging treatment. Furthermore, 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 dissolved back, and with the subsequent primary aging treatment, a γ' phase with a size of 100 to 500 μm can be formed. After the secondary aging treatment, a uniformly dispersed γ' phase with a size of 10 to 50 μm can be formed.

[0094] In a specific embodiment of the present application, the average grain size of the heat-treated alloy is above level 8, with a grain difference of less than level 2. Furthermore, the γ' phase content of the heat-treated alloy is 50 wt% to 58 wt%.

[0095] The nickel-based wrought high-temperature alloy produced by the method of the present application has an average grain size of at least Grade 8, such as Grade 8, Grade 8.5, Grade 9, Grade 9.5, or a range consisting of any two thereof. The grain size difference at various locations in the bar is less than 2 grades, for example, the grain size difference at the center, at R / 2, and at the edge is less than 2 grades, such as Grade 1.5, Grade 1, Grade 0.5, Grade 0, or a range consisting of any two thereof.

[0096] The alloy of the present application ensures a high γ' content, ensuring sufficient strength. For example, in different embodiments, the γ' content can be 50wt%, 51wt%, 52wt%, 53wt%, 54wt%, 55wt%, 56wt%, 57wt%, 58wt%, or a range consisting of any two thereof.

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

[0098] Room temperature tensile strength ≥1510MPa, room temperature yield strength ≥1130MPa;

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

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

[0101] In different embodiments, the tensile properties of the alloy after heat treatment can be exemplified as follows:

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

[0103] The tensile strength at 750°C may be 1110 MPa, 1150 MPa, 1160 MPa, 1180 MPa, 1200 MPa, 1220 MPa, 1240 MPa, 1260 MPa, 1270 MPa, 1275 MPa, or a range composed of any two thereof; the yield strength at 750°C may be 960 MPa, 980 MPa, 1000 MPa, 1020 MPa, 1050 MPa, 1080 MPa, 1090 MPa, or a range composed of any two thereof;

[0104] The tensile strength at 800°C may be 1000 MPa, 1020 MPa, 1050 MPa, 1080 MPa, 1100 MPa, 1120 MPa, 1130 MPa or a range consisting of any two thereof; the yield strength at 800°C may be 900 MPa, 910 MPa, 920 MPa, 930 MPa, 940 MPa, 950 MPa or a range consisting of any two thereof.

[0105] In a specific embodiment of the present application, the endurance life of the alloy after heat treatment satisfies:

[0106] The endurance life of 750℃ / 620MPa is ≥60h;

[0107] The endurance life at 800℃ / 500MPa is ≥45h.

[0108] In different embodiments, the durability of the alloy after heat treatment can be exemplified as follows:

[0109] The endurance life at 750°C / 620 MPa may be 60h, 70h, 80h, 100h, 110h, 120h, 130h, 140h, 150h, 160h, 164h or a range consisting of any two thereof;

[0110] The endurance life at 800°C / 500 MPa may be 45h, 50h, 60h, 70h, 80h, 85h or a range consisting of any two thereof.

[0111] In another aspect, the present application provides a turbine disk made of any one of the above-mentioned nickel-based deformable high-temperature alloys.

[0112] Example 1

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

[0114] A method for preparing a nickel-based deformed high-temperature alloy comprises the following steps:

[0115] (1) According to the alloy composition, a consumable ingot with a specification of Φ508 mm was prepared by a triple smelting process of vacuum induction melting + protective atmosphere electroslag remelting + vacuum consumable remelting, and the surface of the consumable ingot was polished and sprayed with a 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 a low-temperature refining stage at 1480-1510°C. The full melting temperature is controlled at 1520-1550°C, the refining temperature is 1520-1540°C, and the casting temperature is 1440-1450°C.

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

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

[0119] (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; wherein the multi-stage homogenization annealing comprises: 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 1140°C at 50°C / h and holding for 15 hours, then heating to 1180°C at 10°C / h and holding for 25 hours, then heating to 1210°C at 5°C / h and holding for 30 hours, and then furnace cooling.

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

[0121] ① The annealed ingot was kept at 1180℃ for 5h, taken out of the furnace and subjected to upsetting treatment with a pressing amount of 20% and a pressing time of 40s to obtain a primary intermediate billet with a length of 0.8L;

[0122] ② Keep the primary intermediate billet at 1180℃ for 5h, take it out of the furnace and stretch it for 90s to obtain a secondary intermediate billet with a length of L;

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

[0124] ④ Keep the tertiary intermediate billet at 1160℃ for 4h, take it out of the furnace and stretch it for 100s to obtain a quaternary intermediate billet with a length of 1.2L;

[0125] ⑤ Keep the fourth intermediate billet at 1140℃ for 4h, take it out of the furnace and perform upsetting treatment with a pressing amount of 40% and a pressing time of 50s to obtain a fifth intermediate billet with a length of 0.6L;

[0126] ⑥ Keep the fifth intermediate billet at 1140℃ for 4h, take it out of the furnace and stretch it for 120s to obtain a sixth intermediate billet with a size of Φ235mm±10mm;

[0127] ⑦ Preheat the hammer head of the radial forging machine at 500℃ for 4 hours; keep the six intermediate billets at 1130℃ for 5 hours, forge them on the radial forging machine after they are taken out of the furnace, forge them to a size of Φ165mm±10mm, air-cool them to room temperature, and machine them to remove the black skin on the surface to obtain the finished bars with a specification of Φ150mm.

[0128] (4) The finished bar is heat treated; the heat treatment includes: keeping the bar at 1150°C for 4 hours and air cooling to room temperature; then keeping the bar at 860°C for 6 hours and air cooling; and then keeping the bar at 780°C for 24 hours and air cooling to room temperature.

[0129] Example 2

[0130] The present embodiment provides a nickel-based wrought high-temperature alloy, comprising the following components, measured in percentage by mass: 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] A method for preparing a nickel-based deformed high-temperature alloy comprises the following steps:

[0132] (1) According to the alloy composition, a consumable ingot with a specification of Φ508 mm was prepared by a triple smelting process of vacuum induction melting + protective atmosphere electroslag remelting + vacuum consumable remelting, and the surface of the consumable ingot was polished and sprayed with a 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 a low-temperature refining stage at 1480-1510°C. The full melting temperature is controlled at 1540-1550°C, the refining temperature is 1540-1550°C, and the casting temperature is 1440-1450°C.

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

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

[0136] (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; wherein the multi-stage homogenization annealing comprises: 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 1130°C at 50°C / h and holding for 20 hours, then heating to 1170°C at 10°C / h and holding for 30 hours, then heating to 1200°C at 5°C / h and holding for 35 hours, and then furnace cooling.

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

[0138] ① The annealed ingot was kept at 1170℃ for 5h, taken out of the furnace and subjected to upsetting treatment with a pressing amount of 20% and a pressing time of 40s to obtain a primary intermediate billet with a length of 0.8L;

[0139] ② Keep the primary intermediate billet at 1170℃ for 5h, take it out of the furnace and stretch it for 90s to obtain a secondary intermediate billet with a length of L;

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

[0141] ④ Keep the tertiary intermediate billet at 1150℃ for 4h, take it out of the furnace and stretch it for 100s to obtain a quaternary intermediate billet with a length of 1.2L;

[0142] ⑤ Keep the fourth intermediate billet at 1130℃ for 4h, take it out of the furnace and perform upsetting treatment with a pressing amount of 40% and a pressing time of 50s to obtain a fifth intermediate billet with a length of 0.6L;

[0143] ⑥ Keep the fifth intermediate billet at 1130℃ for 4h, take it out of the furnace and stretch it for 120s to obtain the sixth intermediate billet with a size of Φ320mm±10mm;

[0144] ⑦ Preheat the hammer head of the radial forging machine at a temperature of 500°C for 4 hours; keep the six intermediate billets at 1120°C for 5 hours, and forge them on the radial forging machine after they are taken out of the furnace to a size of Φ265mm±10mm, air-cool them to room temperature, and machine them to remove the black skin on the surface to obtain finished bars with a specification of Φ250mm.

[0145] (4) The finished bar is heat treated; the heat treatment includes: keeping the bar at 1150°C for 4 hours and air cooling to room temperature; then keeping the bar at 860°C for 6 hours and air cooling; and then keeping the bar at 780°C for 24 hours and air cooling to room temperature.

[0146] Example 3

[0147] The present embodiment provides a nickel-based wrought high-temperature alloy, comprising the following components, measured in percentage by mass: 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] A method for preparing a nickel-based deformed high-temperature alloy comprises the following steps:

[0149] (1) According to the alloy composition, a consumable ingot with a specification of Φ508 mm was prepared by a triple smelting process of vacuum induction melting + protective atmosphere electroslag remelting + vacuum consumable remelting, and the surface of the consumable ingot was polished and sprayed with a 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 a low-temperature refining stage at 1480-1510°C. The full melting temperature is controlled at 1540-1550°C, the refining temperature is 1540-1550°C, and the casting temperature is 1440-1450°C.

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

[0152] During vacuum consumable remelting, the vacuum consumable melting rate is 3.2kg / 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; wherein the multi-stage homogenization annealing comprises: 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 1120°C at 50°C / h and holding for 25 hours, then heating to 1160°C at 10°C / h and holding for 25 hours, then heating to 1190°C at 5°C / h and holding for 38 hours, and then furnace cooling.

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

[0155] ① The annealed ingot was kept at 1180℃ for 5h, taken out of the furnace and subjected to upsetting treatment with a pressing amount of 20% and a pressing time of 40s to obtain a primary intermediate billet with a length of 0.8L;

[0156] ② Keep the primary intermediate billet at 1180℃ for 4h, take it out of the furnace and stretch it for 90s to obtain a secondary intermediate billet with a length of L;

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

[0158] ④ Keep the tertiary intermediate billet at 1160℃ for 4h, take it out of the furnace and stretch it for 100s to obtain a quaternary intermediate billet with a length of 1.2L;

[0159] ⑤ Keep the fourth intermediate billet at 1140℃ for 4h, take it out of the furnace and perform upsetting treatment with a pressing amount of 40% and a pressing time of 50s to obtain a fifth intermediate billet with a length of 0.6L;

[0160] ⑥ Keep the fifth intermediate billet at 1140℃ for 4h, take it out of the furnace and stretch it for 120s to obtain the sixth intermediate billet with a size of Φ320mm±10mm;

[0161] ⑦ Keep the sixth intermediate billet at 1130℃ for 4h, take it out of the furnace and stretch it for 90s to obtain the seventh intermediate billet with a size of Φ280mm±10mm;

[0162] ⑧Preheat the hammer head of the radial forging machine at 500℃ for 4 hours; keep the six intermediate billets at 1110℃ for 5 hours, forge them on the radial forging machine after taking them out of the furnace, forge them to a size of Φ230mm±10mm, air-cool them to room temperature, and machine them to remove the black skin on the surface to obtain Φ200mm finished bars.

[0163] (4) The finished bar is heat treated; the heat treatment includes: keeping the bar at 1150°C for 4 hours and air cooling to room temperature; then keeping the bar at 860°C for 6 hours and air cooling; and then keeping the bar at 780°C for 24 hours and air cooling to room temperature.

[0164] Example 4

[0165] This embodiment refers to the nickel-based deformed high-temperature alloy and its preparation method of Example 1, with the only difference being that the alloy composition is different.

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

[0167] Example 5

[0168] This embodiment refers to the nickel-based deformed high-temperature alloy and its preparation method of Example 1, with the only difference being that the alloy composition is different.

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

[0170] Example 6

[0171] This embodiment refers to the nickel-based deformed high-temperature alloy and its preparation method of Example 1, with the only difference being that the alloy composition is different.

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

[0173] Example 7

[0174] Example 7 refers to the nickel-based deformed high-temperature alloy and its preparation method of Example 1, with the only difference being that, in the preparation method, in step (2), the homogenization annealing 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 refers to the nickel-based deformed high-temperature alloy and its preparation method of Example 1, with the only difference being that step (3) in the preparation method is different.

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

[0179] ① The annealed ingot was kept at 1180℃ for 5h, taken out of the furnace and subjected to upsetting treatment with a pressing amount of 20% and a pressing time of 40s to obtain a primary intermediate billet with a length of 0.8L;

[0180] ② Keep the primary intermediate billet at 1180℃ for 5h, take it out of the furnace and stretch it for 90s to obtain a secondary intermediate billet with a length of L;

[0181] ③ The secondary intermediate billet is kept at 1180℃ for 4h, taken out of the furnace and subjected to upsetting treatment with a pressing amount of 30% and a pressing time of 45s to obtain a tertiary intermediate billet with a length of 0.7L;

[0182] ④ Keep the tertiary intermediate billet at 1180℃ for 4h, take it out of the furnace and stretch it for 100s to obtain a quaternary intermediate billet with a length of 1.2L;

[0183] ⑤ Keep the fourth intermediate billet at 1180℃ for 4h, take it out of the furnace and perform upsetting treatment with a pressing amount of 40% and a pressing time of 50s to obtain a fifth intermediate billet with a length of 0.6L;

[0184] ⑥ Keep the fifth intermediate billet at 1180℃ for 4h, take it out of the furnace and stretch it for 120s to obtain a sixth intermediate billet with a size of Φ235mm±10mm;

[0185] ⑦ Preheat the hammer head of the radial forging machine at 500℃ for 4 hours; keep the six intermediate billets at 1130℃ for 5 hours, forge them on the radial forging machine after they are taken out of the furnace, forge them to a size of Φ165mm±10mm, air-cool them to room temperature, and machine them to remove the black skin on the surface to obtain the finished bars with a specification of Φ150mm.

[0186] Comparative Example 1

[0187] Comparative Example 1 refers to the nickel-based deformed high-temperature alloy and its preparation method of Example 1, except that the alloy composition is different.

[0188] The nickel-based deformed high-temperature alloy of Comparative Example 1 includes the following components in percentage by mass: 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 deformed high-temperature alloy and its preparation method of Example 1, except that the alloy composition is different.

[0191] The nickel-based deformed high-temperature alloy of Comparative Example 2 includes the following components in percentage by mass: 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 deformed high-temperature alloy and its preparation method of Example 1, except that, in the preparation method, in step (2), the homogenization annealing is different.

[0194] The homogenization annealing in step (2) of comparative example 3 includes: heating to 1210° C. at a rate of 5° C. / h and keeping the temperature for 30 h, and then furnace cooling.

[0195] Comparative Example 4

[0196] Comparative Example 4 refers to the nickel-based deformed high-temperature alloy and its preparation method of Example 2, except that, in the preparation method, step (3) does not include radial forging treatment.

[0197] Step (3) of Comparative Example 4 includes: preparing six intermediate billets according to ① to ⑥ in step (3) of Example 2; then keeping the six intermediate billets at 1120°C for 5 hours, drawing them out of the furnace for 120 seconds, drawing them and rounding them to Φ280mm±10mm, air-cooling them to room temperature, and machining and polishing them to remove the black skin on the surface to obtain finished bars 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. Figures 3 and 4 are grain structure diagrams of the bars prepared in Example 1 and Comparative Example 1 of the present application, respectively; Figure 5 is a microstructure diagram of the annealed alloy ingot prepared in Comparative Example 1; Figures 6 and 7 are 800°C tensile fracture morphology diagrams of the bar prepared in Example 1 of the present application and 800°C / 500 MPa endurance fracture morphology diagrams of the sample; Figures 8 and 9 are 800°C tensile fracture morphology diagrams of the bar prepared in Comparative Example 1 and 800°C / 500 MPa endurance fracture morphology diagrams of the sample.

[0200] Table 1 Metallographic structures of different bars

[0201] Table 2 Mechanical properties test results of different bars

[0202] From the above test results, it can be seen that the nickel-based deformable high-temperature alloy of the present application has the characteristics of high homogeneity and low segregation. The average grain size of the fine-grained rod is finer than level 8, and the grade difference is less than level 2; the room temperature tensile strength is ≥1655MPa, the 750℃ tensile strength is ≥1260MPa, the 750℃ / 620MPa endurance life is ≥145h, and the 800℃ / 500MPa endurance life is ≥72h.

[0203] 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 them. Although the present application has been described in detail with reference to the aforementioned embodiments, those skilled in the art should understand that they can still modify the technical solutions described in the aforementioned embodiments, or make equivalent replacements for some or all of the technical features therein. These modifications or replacements do not deviate the essence of the corresponding technical solutions from the scope of the technical solutions of the embodiments of the present application.

Claims

1. High strength and toughness nickel-based deformable high temperature alloy resistant to 800℃, characterized by: It includes the following components in percentage by mass: 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% to 0.15%, the balance being Ni and unavoidable impurities.

2. The 800°C resistant high-strength and tough nickel-based deformable high-temperature alloy according to claim 1, characterized in that: 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.

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

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

5. The method for preparing the 800°C resistant high-strength and high-toughness nickel-based deformable high-temperature alloy according to any one of claims 1 to 4, characterized in that: The steps include: (a) preparing alloy ingots by a triple smelting process according to the alloy composition; (b) subjecting the alloy ingot to multi-stage homogenization annealing at 1100-1220° C. to obtain an annealed ingot; (c) The annealed ingot is subjected to upsetting and drawing at 1060-1180° C. to form a blank, and then subjected to radial forging at 1020-1150° C.

6. The method for preparing the 800°C resistant high-strength and high-toughness nickel-based deformable high-temperature alloy according to claim 5, characterized in that: The multi-stage homogenization annealing includes: heating the alloy ingot at 490-510° C. for more than 4 hours, then heating it to 990-1010° C. and heating it for more than 8 hours, then heating it to 1110-1140° C. and heating it for more than 15 hours, then heating it to 1150-1180° C. and heating it for more than 25 hours, then heating it to 1190-1220° C. and heating it for more than 30 hours, and then furnace cooling.

7. The method for preparing the 800°C resistant high-strength and high-toughness nickel-based deformable high-temperature alloy according to claim 5, characterized in that: A fast forging machine is used to perform the upsetting and drawing process, wherein the upsetting deformation amount per fire is 10% to 55%, and the drawing deformation amount per fire is 15% to 60%. Preferably, during the upsetting and drawing of the blank, the forging temperature is lowered gradually within the range of 1060-1180°C; Preferably, in the radial forging process, a single-fire multi-pass deformation method is used to prepare an alloy of preset specifications; Preferably, in the radial forging treatment, the holding time at 1020-1150° C. is 30-300 minutes.

8. The method for preparing the 800°C resistant high-strength and tough nickel-based deformable high-temperature alloy according to claim 5, characterized in that: The invention also includes: performing heat treatment on the alloy after the radial forging treatment, wherein the heat treatment includes solution treatment and aging treatment; Preferably, the solution treatment includes: heat preservation treatment at 1100-1150°C for 2-6 hours; the aging treatment includes: heat preservation treatment at 800-900°C for 2-6 hours and then air cooling, and then heat preservation treatment at 750-820°C for 8-24 hours and then air cooling.

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

10. A turbine disc, characterized in that: The alloy is prepared by adopting the 800°C resistant high-strength and tough nickel-based deformable high-temperature alloy described in any one of claims 1 to 4 or the 800°C resistant high-strength and tough nickel-based deformable high-temperature alloy prepared by the preparation method described in any one of claims 5 to 9.

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