Nickel-based powder high-temperature alloy, preparation method thereof and turbine disk

Abstract

The present invention relates to the technical field of high temperature alloys, and in particular to a nickel-based powder high temperature alloy, a preparation method thereof and a turbine disk. The nickel-based powder high temperature alloy comprises C 0.05% to 0.10%, Cr 10.5% to 13.5%, Co 13% to 16%, Mo 4% to 6%, W 2.9% to 4.1%, Al 3.9% to 4.5%, Ti 2.7% to 3.3%, Nb 3.2% to 3.8%, Hf 0.01% to 0.15%, V 0.5% to 1.0%, La 0.01% to 0.04%, Zr 0.01% to 0.04%, Mg 0.005% to 0.015%, B 0.005% to 0.03%, Ce 0.005% to 0.015%, and the balance is Ni. The present invention adjusts and controls the alloy composition and matches the corresponding heat treatment conditions to solve the problem of insufficient strength margin of existing high-temperature alloys without substantially losing plasticity.

Description

Technical Field

[0001] The invention relates to the technical field of high-temperature alloys, and in particular to a nickel-based powder high-temperature alloy and a preparation method thereof and a turbine disk. Background Art

[0002] The turbine disk is an important core hot end component of an aircraft engine. Its metallurgical quality and performance level play a decisive role in improving the reliability, safety life and performance of the engine and aircraft. As a precipitation-strengthened powder high-temperature alloy, FGH4097 alloy has excellent comprehensive mechanical properties and is a key material for key hot end components such as aircraft engine turbine disks. At present, the FGH4097 alloy powder turbine disk formed by direct hot isostatic pressing can meet the requirements of the full life target of a certain engine. With the emergence of new engines with higher performance and higher requirements, the strength margin of FGH4097 alloy is insufficient, and the service life margin of the turbine disk is limited, which cannot meet the requirements of higher performance engines.

[0003] In view of this, the present invention is proposed. Summary of the invention

[0004] The purpose of the present invention is to provide a nickel-based powder high-temperature alloy, a preparation method thereof and a turbine disk. The present invention significantly improves the strength margin of the nickel-based powder high-temperature alloy by reasonably adjusting the alloy composition, and has better high-temperature endurance strength and low-cycle fatigue strength, which can significantly improve the service life margin of the turbine disk.

[0005] In order to achieve the above-mentioned purpose of the present invention, the first aspect of the present invention provides a nickel-based powder high-temperature alloy, which includes, by mass percentage, C 0.05%-0.10%, Cr 10.5%-13.5%, Co 13%-16%, Mo 4%-6%, W 2.9%-4.1%, Al 3.9%-4.5%, Ti 2.7%-3.3%, Nb 3.2%-3.8%, Hf 0.01%-0.15%, V 0.5%-1.0%, La 0.01%-0.04%, Zr 0.01%-0.04%, Mg 0.005%-0.015%, B 0.005%-0.03%, Ce0.005%-0.015%, and the balance is Ni and unavoidable impurities.

[0006] In a specific embodiment of the present invention, in the nickel-based powder high-temperature alloy, by mass percentage, O≤0.015%, Fe≤1%, Mn≤0.3%, Si≤0.3%, S≤0.009%, P≤0.015%, and N≤0.001%.

[0007] In a specific embodiment of the present invention, in the nickel-based powder high-temperature alloy, the ratio of the mass percentage of Al to the mass percentage of Nb, Al / Nb, is 1.1 to 1.35.

[0008] In a specific embodiment of the present invention, in the nickel-based powder high-temperature alloy, the ratio of the mass percentage of Al to the mass percentage of Ti, Al / Ti, is 1.3 to 1.6.

[0009] In a specific embodiment of the present invention, in the nickel-based powder high-temperature alloy, the ratio of the mass percentage of W to the mass percentage of Mo, W / Mo, is 0.6 to 0.8.

[0010] In a specific embodiment of the present invention, in the nickel-based powder high-temperature alloy, the ratio of the mass percentage of Hf to the mass percentage of C (Hf / C) is ≤1.5.

[0011] The second aspect of the present invention provides a method for preparing the nickel-based powder high-temperature alloy of the first aspect, comprising the following steps:

[0012] (a) preparing a master alloy bar by vacuum induction melting according to the alloy composition;

[0013] (b) preparing alloy powder from the master alloy rod by a plasma rotating electrode method;

[0014] (c) hot isostatic pressing the alloy powder to obtain a blank;

[0015] (d) subjecting the blank to heat treatment.

[0016] In a specific embodiment of the present invention, in the hot isostatic pressing forming, the hot isostatic pressing temperature is 1160-1200°C.

[0017] In a specific embodiment of the present invention, the heat treatment includes: heat preservation treatment at 1150-1200°C for 1-12 hours, air cooling, and then two-stage aging treatment. Further, the two-stage aging treatment includes: heat preservation treatment at 750-800°C for 6-20 hours, air cooling, and then heat preservation treatment at 650-740°C for 6-20 hours, and then air cooling.

[0018] The third aspect of the present invention provides a turbine disk, which is mainly made of the nickel-based powder high-temperature alloy described in the first aspect.

[0019] Compared with the prior art, the present invention has the following beneficial effects:

[0020] (1) The present invention regulates the alloy composition, increases the contents of C, Cr, Mo, Ti, and Nb, reduces the contents of W, Co, and Al, and introduces a certain amount of La element to increase the solid solution temperature of the γ′ strengthening phase, thereby ensuring a wider single-phase temperature range. In combination with corresponding heat treatment conditions, the problem of insufficient strength margin of existing high-temperature alloys is solved without substantially losing plasticity.

[0021] (2) The high-temperature alloy of the present invention can be directly formed by hot isostatic pressing, which has the advantages of short process and low cost, while ensuring the performance of the formed parts;

[0022] (3) The alloy of the present invention has excellent tensile strength, 650°C endurance strength and low-cycle fatigue strength, which can significantly improve the service life of the turbine disk and meet the material selection requirements of new engines with higher performance. BRIEF DESCRIPTION OF THE DRAWINGS

[0023] In order to more clearly illustrate the specific implementation methods of the present invention or the technical solutions in the prior art, the drawings required for use in the specific implementation methods or the description of the prior art will be briefly introduced below. Obviously, the drawings described below are some implementation methods of the present invention. For ordinary technicians in this field, other drawings can be obtained based on these drawings without paying creative work.

[0024] Figure 1 The calculated phase diagram of the nickel-based powder high-temperature alloy provided in Example 1 of the present invention;

[0025] Figure 2 The calculated phase diagram of the FGH4097 alloy provided in Comparative Example 1 of the present invention;

[0026] Figure 3 This is an electron microscope image of the FGH4097 alloy provided in Comparative Example 1 of the present invention;

[0027] Figure 4 The metallographic structure of the hot isostatically pressed alloy provided in Example 1 of the present invention;

[0028] Figure 5 The metallographic structure of the hot isostatically pressed alloy provided in Example 7 of the present invention;

[0029] Figure 6 The metallographic structure of the hot isostatically pressed alloy provided in Example 8 of the present invention;

[0030] Figure 7 The metallographic structure of the heat-treated alloy provided in Example 1 of the present invention;

[0031] Figure 8 This is an electron microscope image of the heat-treated alloy provided in Example 1 of the present invention. DETAILED DESCRIPTION

[0032] The technical scheme of the present invention 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 invention, rather than all embodiments, and are only used to illustrate the present invention, and should not be considered as limiting the scope of the present invention. Based on the embodiments in the present invention, 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 invention. If specific conditions are not specified in the embodiments, they are carried out according to normal conditions or conditions recommended by the manufacturer. If the manufacturer is not specified in the reagents or instruments used, they are all conventional products that can be purchased commercially.

[0033] With the emergence of new engines with higher performance and higher requirements, the strength margin of FGH4097 alloy is insufficient. Given the significant advantages of direct hot isostatic pressing technology, which has a short process and low cost, obtaining powder high-temperature alloys with significantly improved strength margin by regulating the alloy composition under the premise of adopting hot isostatic pressing technology is an inevitable requirement for improving the service life of turbine disks, and it is also a technical problem in this field.

[0034] Based on this, the first aspect of the present invention provides a nickel-based powder high-temperature alloy, which includes, by mass percentage, C 0.05%~0.10%, Cr 10.5%~13.5%, Co 13%~16%, Mo 4%~6%, W 2.9%~34.1%, Al 3.9%~4.5%, Ti 2.7%~3.3%, Nb 3.2%~3.8%, Hf 0.01%~0.15%, V 0.5%~1.0%, La 0.01%~0.04%, Zr 0.01%~0.04%, Mg 0.005%~0.015%, B 0.005%~0.03%, Ce 0.005%~0.015%, and the balance is Ni and unavoidable impurities.

[0035] Compared with the FGH4097 alloy, the nickel-based powder high-temperature alloy of the present invention has higher tensile strength, 650°C endurance strength and low-cycle fatigue strength without losing plasticity, thereby solving the problem of insufficient strength margin of existing high-temperature alloys and meeting the material selection requirements of new engines with higher performance.

[0036] Compared with the FGH4097 alloy, the present invention changes and regulates the composition of the high-temperature alloy, increases the contents of C, Cr, Mo, Ti, Nb, etc., reduces the contents of W, Co, Al, Hf, etc., and introduces a certain amount of La element. On the premise that the total amount of elements forming the γ′ strengthening phase remains basically unchanged, the solid solution temperature of the γ′ strengthening phase is increased, a wider single-phase zone temperature is ensured, and corresponding heat treatment conditions are combined to solve the problem of insufficient strength margin of existing high-temperature alloys.

[0037] For example, in different embodiments, the content of C element in the nickel-based powder high-temperature alloy can be 0.05%, 0.06%, 0.07%, 0.08%, 0.09%, 0.10% or a range consisting of any two thereof, calculated by mass percentage.

[0038] For example, in different embodiments, the content of Cr in the nickel-based powder high-temperature alloy may be 10.5%, 11%, 11.5%, 12%, 12.5%, 13%, 13.5% or a range consisting of any two thereof, calculated by mass percentage.

[0039] For example, in different embodiments, the content of Co in the nickel-based powder high-temperature alloy can be 13%, 13.5%, 14%, 14.5%, 15%, 15.5%, 16% or a range consisting of any two thereof, calculated by mass percentage.

[0040] For example, in different embodiments, the content of Mo element in the nickel-based powder high-temperature alloy can be 4%, 4.2%, 4.5%, 4.8%, 5%, 5.2%, 5.5%, 5.8%, 6% or a range consisting of any two thereof, calculated by mass percentage.

[0041] For example, in different embodiments, the content of W in the nickel-based powder high-temperature alloy can be 2.9%, 3%, 3.2%, 3.5%, 3.8%, 4%, 4.1% or any two of them in terms of mass percentage. Studies have found that in the alloy system of the present invention, an appropriate amount of W content can improve the stability of the γ / γ′ phase interface and the stability of the structure, increase the activation energy of atomic diffusion, slow down the atomic diffusion process, increase the recrystallization temperature, and significantly reduce the steady-state creep rate. However, W is also a typical TCP phase-forming element, which will reduce the energy required for the formation of μ phase; the present invention regulates the content of W within the above range to avoid it inducing the precipitation of μ phase.

[0042] For example, in different embodiments, the content of Al element in the nickel-based powder high-temperature alloy can be 3.9%, 4%, 4.1%, 4.2%, 4.4%, 4.5% or a range consisting of any two thereof, calculated by mass percentage.

[0043] For example, in different embodiments, the content of Ti element in the nickel-based powder high-temperature alloy can be 2.7%, 2.8%, 3%, 3.1%, 3.2%, 3.3% or a range consisting of any two thereof, calculated by mass percentage.

[0044] For example, in different embodiments, the content of Nb in the nickel-based powder high-temperature alloy can be 3.2%, 3.3%, 3.4%, 3.5%, 3.6%, 3.8% or a range composed of any two of them, measured by mass percentage; the alloy of the present invention adopts a higher Nb content, which can refine the γ′ phase and improve the performance of the alloy, and the higher content of Nb can partially replace the Al site in the lattice, thereby increasing the anti-phase boundary energy of the γ′ phase and improving the dispersion strengthening effect of the γ′ phase in the γ matrix; at the same time, the higher content of Nb can be in the Cr 2 O 3 A continuous Nb-rich layer is formed under the oxide layer, which prevents Ni from diffusing outward during oxidation and improves the oxidation resistance of the alloy. When the Nb content is too high, the continuous Nb-rich layer will be converted into a discontinuous layer. During the oxidation process, O can more easily penetrate the discontinuous layer into the alloy matrix, thereby greatly reducing the oxidation resistance of the alloy.

[0045] For example, in different embodiments, the content of Hf in the nickel-based powder high-temperature alloy can be 0.01%, 0.03%, 0.05%, 0.06%, 0.07%, 0.08%, 0.09%, 0.1%, 0.12%, 0.15% or a range consisting of any two of them, by mass percentage. Hf is a strong MC carbide-forming element. The addition of Hf can weaken carbon segregation on grain boundaries and inhibit the degradation of MC carbides, thereby preventing M 23 C 6 The formation of microtwins can be slowed down by Hf, which can slow down the initiation and expansion of cracks, thereby improving the creep properties of the alloy; at the same time, Hf can increase the content and equivalent particle size of secondary γ′, increase the stacking fault energy of γ′ phase, and thus inhibit the expansion of microtwins. However, excessive Hf will directly form a large amount of HfC and HfO 2 And it will enter the MC' type carbide, which is not conducive to the effective replacement of Nb, Ti and other elements in the MC phase by the Hf element. At the same time, excessive Hf elements will significantly reduce the solidus temperature and increase the complete dissolution temperature of the γ' phase, which is not conducive to the alloy solution heat treatment process. Therefore, the present invention regulates the Hf element content within the above range.

[0046] For example, in different embodiments, the content of V element in the nickel-based powder high-temperature alloy can be 0.5%, 0.6%, 0.7%, 0.8%, 0.9%, 1.0% or a range consisting of any two thereof, calculated by mass percentage.

[0047] For example, in different embodiments, the content of La in the nickel-based powder high-temperature alloy can be 0.01%, 0.015%, 0.02%, 0.025%, 0.03%, 0.035%, 0.04% or any two of them, in terms of mass percentage. The addition of La can combine with S to form a compound, reduce S in the alloy, and purify the grain boundaries; at the same time, La can hinder atomic diffusion, thereby inhibiting M 23 C 6 growth and thinning, while the small M 23 C 6 It can effectively hinder dislocation movement, thereby improving the strength and creep properties of the alloy. At the same time, La segregates at the oxide-metal interface, inhibiting Mn 2+ and Cr 3+ The diffusion rate of the plasma is reduced, thereby slowing down the oxidation process and improving the oxidation resistance of the alloy. However, as the La content increases, the oxidation rate also increases, mainly because a large amount of La-Ni phase is oxidized to form La-O oxides, which accelerates the oxidation process. Therefore, the present invention regulates the La element content within the above range.

[0048] For example, in different embodiments, the content of Zr in the nickel-based powder high-temperature alloy can be 0.01%, 0.015%, 0.02%, 0.025%, 0.03%, 0.035%, 0.04% or any two thereof, by mass percentage; by regulating the content of Zr within the above range, on the one hand, Zr can promote the refinement and spheroidization of MC-type carbides and effectively prevent grain boundary sliding; on the other hand, Zr can also significantly increase the binding energy between the precipitate phase / matrix interface and the grain boundary, which is beneficial to stress regulation and delays crack initiation and propagation; in the process of long-term high-temperature aging or high-temperature creep, Zr helps to improve the thermal stability of MC and inhibit its transformation to M 23 C 6 The Zr element is concentrated at the grain boundary, which improves the structural integrity and grain boundary bonding strength, reduces the interface energy between carbides and the matrix, and helps to hinder dislocation slip and crack propagation; in addition, the Zr element can also reduce the precipitation rate of carbides at the grain boundary and the γ′ depletion rate near the grain boundary, delaying the microstructural process that induces microcracks. At the same time, the Zr element has low solubility and strong positive segregation in the γ matrix of the Ni-based high-temperature alloy. The present invention further regulates the Zr element content within the above range to avoid excessive Zr elements causing the low melting point phase rich in Zr elements to appear in the alloy, resulting in the initial melting phenomenon during the hot working process.

[0049] For example, in different embodiments, in the nickel-based powder high-temperature alloy, the content of Mg element can be 0.005%, 0.008%, 0.01%, 0.012%, 0.015% or a range consisting of any two thereof, by mass percentage; the content of B element can be 0.005%, 0.01%, 0.015%, 0.02%, 0.025%, 0.03% or a range consisting of any two thereof; the content of Ce element can be 0.005%, 0.008%, 0.01%, 0.012%, 0.015% or a range consisting of any two thereof.

[0050] In a specific embodiment of the present invention, in the nickel-based powder high-temperature alloy, by mass percentage, O≤0.015%, Fe≤1%, Mn≤0.3%, Si≤0.3%, S≤0.009%, P≤0.015%, and N≤0.001%.

[0051] In the nickel-based powder high-temperature alloy of the present invention, the content of each impurity can be controlled within the above range. The content of O in the nickel-based powder high-temperature alloy of the present invention can be controlled to be below 0.015%, while the content of O in the existing FGH4097 needs to be controlled to be below 0.007%. Compared with FGH4097, the content of O in the alloy of the present invention can be relatively increased, which significantly reduces the production cost of the alloy.

[0052] In a specific embodiment of the present invention, in the nickel-based powder high-temperature alloy, the ratio of the mass percentage of Al to the mass percentage of Nb, Al / Nb, is 1.1 to 1.35, for example, it can be 1.1, 1.15, 1.2, 1.25, 1.3, 1.35 or a range composed of any two of them. Al is one of the main elements for the formation of the γ' phase. For the alloy system of the present invention, when Al / Nb is too low, sufficient γ' phase cannot be formed in the alloy, which will lead to a decrease in the strength of the alloy; when Al / Nb is too high, Nb cannot replace the point of Al in the lattice, thereby failing to increase the antiphase domain boundary energy of the γ' phase, and the strength of the alloy will also decrease. The present invention regulates Al / Nb within the above range, which is more conducive to improving the strength of the alloy.

[0053] In a specific embodiment of the present invention, in the nickel-based powder high-temperature alloy, the ratio of the mass percentage of Al to the mass percentage of Ti, Al / Ti, is 1.3 to 1.6, for example, it can be 1.3, 1.35, 1.4, 1.45, 1.5, 1.55, 1.6 or a range composed of any two of them, which helps to further increase the solution temperature of the γ' strengthening phase while ensuring the amount of the γ' phase, thereby ensuring a wider single-phase region temperature.

[0054] In a specific embodiment of the present invention, in the nickel-based powder high-temperature alloy, the ratio of the mass percentage of W to the mass percentage of Mo, W / Mo, is 0.6 to 0.8, for example, it can be 0.6, 0.65, 0.7, 0.75, 0.8 or a range composed of any two thereof, thereby further improving the strength and creep properties of the alloy.

[0055] In a specific embodiment of the present invention, in the nickel-based powder high-temperature alloy, the ratio of the mass percentage of Hf to the mass percentage of C is Hf / C≤1.5, for example, it can be 1.5, 1.2, 1, 0.8, 0.7 or a range composed of any two of them, thereby giving full play to the improvement of the creep properties of the alloy.

[0056] The second aspect of the present invention provides a method for preparing the nickel-based powder high-temperature alloy of the first aspect, comprising the following steps:

[0057] (a) preparing a master alloy bar by vacuum induction melting according to the alloy composition;

[0058] (b) preparing alloy powder from master alloy rod by plasma rotating electrode method;

[0059] (c) hot isostatic pressing the alloy powder to obtain a blank;

[0060] (d) subjecting the blank to heat treatment.

[0061] In a specific embodiment of the present invention, the particle size of the alloy powder is less than 100 μm. The present invention can adopt a conventional plasma rotating electrode method to prepare the alloy powder, and ensure that the particle size of the alloy powder is within the above range. Furthermore, the number of ceramics and other inclusions in each 1.0 kg of finished powder is less than 20, and the maximum inclusion size is not greater than 100 μm.

[0062] In a specific embodiment of the present invention, in the hot isostatic pressing forming, the temperature of the hot isostatic pressing is 1160-1200°C, for example, it can be 1160°C, 1170°C, 1180°C, 1190°C, 1200°C or a range composed of any two thereof. The present invention has found that when the temperature of the hot isostatic pressing is within 1180-1190°C, the hot isostatic pressed alloy obtained has a suitable grain size of 7.0-7.5, which is more conducive to improving the high temperature performance while ensuring the room temperature tensile performance. When the temperature of the hot isostatic pressing is lower than 1180°C, the grain size of the hot isostatic pressed alloy is 8.0-8.5, the dendrites in the organization are completely eliminated, the grain structure is relatively uniform, and the room temperature tensile performance is not significantly different from the hot isostatic pressed alloy at 1180-1190°C, but the high temperature tensile performance at 650°C is lower than that of the hot isostatic pressed alloy at 1180-1190°C. This may be because under high temperature conditions of 650℃, the grain boundary strength is lower than the intragranular strength, and crack sources are more likely to initiate and expand at the grain boundaries, while finer grain sizes mean more grain boundaries, which leads to a sharp drop in the performance of the alloy at high temperatures. When the hot isostatic pressing temperature is higher than 1190℃, the grain size of the hot isostatically pressed alloy is 6.3 to 7.2, and the uniformity of the grain structure is worse than that of the hot isostatically pressed alloy at 1180 to 1190℃.

[0063] In a specific embodiment of the present invention, in hot isostatic pressing, the hot isostatic pressing pressure is 140-180 MPa, for example, it can be 140 MPa, 150 MPa, 160 MPa, 170 MPa, 180 MPa or a range composed of any two of them, and the holding time is 2-6 h, for example, it can be 2 h, 4 h, 6 h, etc.

[0064] In a specific embodiment of the present invention, the grain size of the blank is 7.0-7.5.

[0065] In a specific embodiment of the present invention, the heat treatment includes: heat preservation treatment at 1150-1200°C for 1-12 hours, air cooling, and then two-stage aging treatment. Further, the two-stage aging treatment includes: heat preservation treatment at 750-800°C for 6-20 hours, air cooling, and then heat preservation treatment at 650-740°C for 6-20 hours, and then air cooling.

[0066] For example, in different embodiments, the temperature of the solution treatment can be 1150°C, 1160°C, 1170°C, 1180°C, 1190°C, 1200°C, or a range consisting of any two of them. After the solution treatment at a temperature of 1180-1190°C, the grain size of the heat-treated alloy is 7.5-8.0, no PPBs (Prior Particle Boundarys) defects are found in the structure, and no obvious dendritic structure appears. When the temperature of the solution treatment is lower than 1180°C, there will be large-sized γ′ that is not dissolved, thereby reducing the number of secondary and tertiary γ′ phases precipitated during subsequent cooling, thereby reducing the strengthening effect of the γ′ phase. When the temperature of the solution treatment is higher than 1190°C, the primary γ′ phase at the grain boundary dissolves, causing the grain boundary to lose its pinning effect, thereby causing the grain to grow and reducing the strength of the alloy. The heat-treated state of the present invention refers to after solution and aging treatment.

[0067] For example, in different embodiments, in the two-stage aging treatment, the temperature of the primary aging can be 750°C, 760°C, 770°C, 780°C, 790°C, 800°C or a range consisting of any two thereof, and the temperature of the primary aging can be 6h, 8h, 12h, 14h, 15h, 16h, 20h or a range consisting of any two thereof. When the primary aging is regulated to meet the above conditions, it is helpful to fully precipitate the tertiary γ′ phase and evenly distribute it in the γ channel, greatly improving the strength of the alloy, and significantly improving the tensile strength and yield strength of the alloy at a high temperature of 650°C. Studies have found that when the temperature of the primary aging is higher than 760-780°C, or the time of the primary aging is shorter than the above range, it may cause the tertiary γ′ phase to be unable to be fully precipitated, or it may cause the tertiary γ′ phase to grow further, lose the strengthening effect or weaken the strengthening effect, which will lead to a decrease in the tensile properties of the alloy.

[0068] The third aspect of the present invention provides a turbine disk, which is mainly made of the nickel-based powder high-temperature alloy of the first aspect.

[0069] In a specific embodiment of the present invention, the tensile strength of the turbine disk at 750° C. is ≥1210 MPa, the yield strength is ≥1000 MPa, and the elongation is ≥20%.

[0070] Example 1

[0071] The present embodiment provides a nickel-based powder high-temperature alloy, including the following components in mass percentage: C0.07%, Cr 12.0%, Co 14.5%, Mo 5.0%, W 3.5%, Al 4.2%, Ti 3.0%, Nb 3.5%, Hf 0.075%, V0.75%, La 0.03%, Zr 0.025%, Mg 0.015%, B 0.025%, Ce 0.015%, O≤0.015%, Fe≤1%, Mn≤0.3%, Si≤0.3%, S≤0.009%, P≤0.015%, N≤0.001% and Ni balance.

[0072] A method for preparing a nickel-based powder high-temperature alloy comprises the following steps:

[0073] (1) The master alloy bar with a specification of φ63×700 mm is prepared according to the alloy composition and vacuum induction melting; wherein, vacuum induction melting includes: all raw materials are used after being baked, and materials are added according to the alloy composition. After the alloy liquid is melted and refined, chemical sampling and testing are carried out before the furnace; according to the composition results before the furnace and the internal control requirements of the elements, additional materials are added to fine-tune the composition, and the melting is completed by temperature measurement and pouring.

[0074] (2) The master alloy rod obtained in step (1) is made into alloy powder by a plasma rotating electrode method, and a finished powder with a particle size of less than 100 μm is collected, and the number of ceramics and other inclusions per 1.0 kg of finished powder is less than 20, and the maximum inclusion size is not greater than 90 μm; wherein the plasma rotating electrode method includes: the rod is φ60×700 mm, the rotation speed is 1.5 w rpm, and the prepared powder particle size is less than 100 μm.

[0075] (3) The alloy powder obtained in step (2) is placed in a package, vacuum degassed and sealed, and then hot isostatic pressing is performed to obtain a blank; wherein the hot isostatic pressing temperature is 1185°C, the pressure is 160 MPa, and the time is 6 hours.

[0076] (4) heat treating the blank obtained in step (3); wherein the heat treatment comprises: heat treatment at 1185°C for 2 hours and then air cooling; then heat treatment at 770°C for 14 hours and then air cooling; then heat treatment at 700°C for 14 hours and then air cooling.

[0077] Embodiments 2 to 6

[0078] Examples 2 to 6 refer to the nickel-based powder high temperature alloy and its preparation method of Example 1, and the only difference is that the alloy component content is different. The alloy component content of Examples 2 to 6 is shown in Table 1.

[0079] Table 1 Alloy composition content of Examples 2 to 6

[0080]

[0081] Embodiment 7-8

[0082] Example 7 and Example 8 refer to the nickel-based powder high-temperature alloy and its preparation method of Example 1 respectively, and the only difference is that in the preparation method, the temperature of the hot isostatic pressing in step (3) is different.

[0083] In Example 7, the temperature of hot isostatic pressing is 1170°C;

[0084] In Example 8, the temperature of hot isostatic pressing is 1200°C.

[0085] Embodiments 9-10

[0086] Example 9 and Example 10 refer to the nickel-based powder high-temperature alloy and its preparation method of Example 1 respectively, and the only difference is that in the preparation method, in the heat treatment of step (4), the two-stage aging treatment system is different.

[0087] In Example 9, the two-stage aging treatment includes: heat preservation treatment at 770°C for 8 hours and then air cooling, and then heat preservation treatment at 700°C for 14 hours and then air cooling;

[0088] In Example 10, the two-stage aging treatment includes: heat preservation treatment at 790° C. for 14 hours and then air cooling, and then heat preservation treatment at 700° C. for 14 hours and then air cooling.

[0089] Comparative Example 1

[0090] Comparative Example 1 provides a FGH4097 alloy comprising the following components in mass percentage: C 0.04%, Cr 9%, Co 16%, Mo 3.9%, W 5.6%, Al 5.2%, Ti 1.9%, Nb 2.7%, Hf 0.25%, Zr≤0.015%, Mg≤0.02%, B≤0.03%, Ce≤0.01%, O≤0.007% and Ni balance.

[0091] The preparation method of FGH4097 alloy includes: master alloy smelting, PREP powder making, jacket degassing and sealing welding, hot isostatic pressing, and then heat treatment. Among them, the hot isostatic pressing temperature is 1150-1200℃ (such as 1185℃), the time is 6h, and the pressure is greater than 140MPa (such as 160MPa); in the heat treatment, the solid solution temperature is 1140-1200℃ (such as 1185℃), the time is 2-8h (such as 2h), the aging temperature is 700-900℃ (such as 800℃), and the time is 15-30h (such as 16h).

[0092] Comparative Examples 2 to 4

[0093] Comparative Examples 2 to 4 refer to the nickel-based powder high temperature alloy and its preparation method of Example 1, except that the contents of alloy components are different. The contents of alloy components of Comparative Examples 2 to 4 are shown in Table 2.

[0094] Table 2 Alloy composition content of comparative examples 2 to 4

[0095]

[0096] Experimental Example 1

[0097] Figure 1 and Figure 2 The calculations are the same for the nickel-based powder high-temperature alloy of Example 1 of the present invention and the FGH4097 alloy of Comparative Example 1. As can be seen from the figure, M 6 C carbide, but M does not appear in the alloy phase diagram calculation of Example 1 of the present invention. 6 C carbide, this is because the W content in the alloy of Comparative Example 1 is much higher than that in the alloy of Example 1 of the present invention, and W is M 6 C is the main element in carbide formation, so M will appear in the alloy of comparative example 1. 6 C carbide, and when M 6 When C carbides form a continuous film, it will seriously affect the durability of the alloy. Figure 3 This is an electron microscope image of the alloy obtained in Comparative Example 1. A large amount of M 6 C carbide.

[0098] Experimental Example 2

[0099] The metallographic structures of the hot isostatically pressed alloys provided in Examples 1 and 7-8 of the present invention are characterized, respectively. Figures 4 to 6 As shown in the figure. It can be seen that the grain size of the alloy obtained by the hot isostatic pressing temperature of Example 1 of the present invention is 7.0-7.5; the grain size of the alloy obtained by the hot isostatic pressing temperature of Example 7 of the present invention is 8.0-8.5; the grain size of the alloy obtained by the hot isostatic pressing temperature of Example 8 of the present invention is 6.3-7.2. In addition, the dendrites in the hot isostatically pressed alloy obtained in Example 7 of the present invention have been completely eliminated, and the grain structure is relatively uniform.

[0100] The high temperature tensile properties of the hot isostatically pressed alloys of Example 1 and Example 7 of the present invention were further tested. The test results are shown in Table 3.

[0101] Table 3 High temperature tensile properties test results of the alloys of Example 1 and Example 7

[0102]

[0103] From the above test results, it can be seen that the 650°C high temperature tensile property of the hot isostatically pressed alloy of Example 7 of the present invention is lower than the 650°C high temperature tensile property of the hot isostatically pressed alloy of Example 1 of the present invention. The reason may be that at a high temperature of 650°C, the grain boundary strength is lower than the intragranular strength, and the crack source is more likely to initiate and expand at the grain boundary, and a finer grain size means more grain boundaries, which leads to a decrease in the performance of the alloy of Example 7 at high temperatures.

[0104] The metallographic structure of the heat-treated alloy provided in Example 1 of the present invention is further characterized. Figure 7-Figure 8 The alloy obtained in Example 1 of the present invention has a grain size of 7.5 to 8.0, an average grain size of 22.5 to 26.7 μm, no PPBs defects are found, no obvious dendrite structure appears, and the average size of the γ′ phase is about 0.3 μm.

[0105] The tensile properties of the heat-treated alloys provided in Examples 1 to 6, Examples 9 to 10 and Comparative Examples 1 to 4 of the present invention were further tested, and the test results are shown in Tables 4 to 6.

[0106] Table 4 Test results of room temperature tensile properties of different alloys

[0107]

[0108] Table 5 650℃ tensile properties test results of different alloys

[0109]

[0110] Table 6 Tensile properties test results of different alloys at 750℃

[0111]

[0112] The composite endurance performance and low cycle fatigue performance of the alloy prepared in Example 1 of the present invention and the alloy prepared in Comparative Example 1 were further tested, and the test results are shown in Tables 7 and 8.

[0113] Table 7 Composite durability performance

[0114]

[0115] Table 8 Low cycle fatigue properties

[0116]

[0117] It can be seen from the above test results that, compared with the FGH4097 alloy, the tensile properties, endurance properties and low-cycle fatigue properties of the alloy of the present invention are significantly improved, which can significantly increase the service life margin of the turbine disk blank.

[0118] Finally, it should be noted that the above embodiments are only used to illustrate the technical solutions of the present invention, rather than to limit it. Although the present invention 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 replace some or all of the technical features therein with equivalents. 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 invention.

Claims

1. Nickel-based powder high temperature alloy, characterized in that: In terms of mass percentage, it includes: C 0.05%-0.10%, Cr 10.5%-13.5%, Co 13%-16%, Mo 4%-6%, W 2.9%-4.1%, Al 3.9%-4.5%, Ti 2.7%-3.3%, Nb3.2%-3.8%, Hf 0.01%-0.15%, V 0.5%-1.0%, La 0.01%-0.04%, Zr 0.01%-0.04%, Mg0.005%-0.015%, B 0.005%-0.03%, Ce 0.005%-0.015%, and the balance is Ni and unavoidable impurities; The ratio of the mass percentage of Al to the mass percentage of Nb, Al / Nb, is 1.1 to 1.35; The ratio of the mass percentage of Al to the mass percentage of Ti, Al / Ti, is 1.3 to 1.6; The ratio of the mass percentage of W to the mass percentage of Mo, W / Mo, is 0.6 to 0.8; The ratio of the mass percentage of Hf to the mass percentage of C is Hf / C ≤ 1.5; A method for preparing a nickel-based powder high-temperature alloy comprises the following steps: (a) preparing a master alloy rod by vacuum induction melting according to the composition of the alloy; (b) preparing alloy powder from the master alloy rod by a plasma rotating electrode method; (c) hot isostatic pressing the alloy powder to obtain a blank; (d) subjecting the blank to heat treatment; The particle size of the alloy powder is less than 100 μm; In the hot isostatic pressing, the temperature of the hot isostatic pressing is 1160-1190°C; The heat treatment comprises: heat preservation treatment at 1150-1200°C for 1-12 hours, air cooling, and then two-stage aging treatment; The two-stage aging treatment includes: heat preservation treatment at 750-800°C for 6-20 hours and then air cooling, and then heat preservation treatment at 650-700°C for 6-20 hours and then air cooling.

2. The nickel-based powder high temperature alloy according to claim 1, characterized in that: In terms of mass percentage, O≤0.015%, Fe≤1%, Mn≤0.3%, Si≤0.3%, S≤0.009%, P≤0.015%, and N≤0.001%.

3. A turbine disc, characterized in that: The invention is mainly made of the nickel-based powder high-temperature alloy as claimed in claim 1 or 2.

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