A nickel-based superalloy and a method of making the same

Abstract

The application discloses a nickel-based high-temperature alloy and a preparation method thereof, and belongs to the technical field of high-temperature alloys. The nickel-based high-temperature alloy can meet the requirements of the service environment of 700 DEG C to 750 DEG C and hot workability of a hot end component of an aero-engine. The components of the nickel-based high-temperature alloy include the following components in percentage by mass: Mg: 0.01% to 0.02%, Cu: 0.005% to 0.1%, Fe: 3.8% to 4.8%, Nb: 1% to 1.3%, Cr: 15.2% to 16.5%, Co: 9% to 9.98%, Mo: 2.7% to 3.4%, W: 2.5% to 2.9%, Al: 2.1% to 2.5%, Ti: 3.3% to 3.7%, Zr: 0.025% to 0.05%, C: 0.01% to 0.014%, B: 0.005% to 0.014%, and nickel: the rest. The nickel-based high-temperature alloy has excellent high-temperature mechanical properties and hot workability.

Description

Technical Field

[0001] This invention relates to the field of high-temperature alloy technology, and in particular to a nickel-based high-temperature alloy and its preparation method. Background Technology

[0002] Nickel-based superalloys have become core materials for the preparation of hot-end components of aero-engines due to their excellent high-temperature strength, oxidation resistance and long-term structural stability, and are widely used in high-end equipment fields such as spacecraft, marine gas turbines and nuclear power equipment.

[0003] Currently, general-purpose high-temperature alloys, represented by GH4169 alloy, are the core backbone materials with the largest usage and widest coverage in my country's strategic fields such as aviation, aerospace, and energy. However, with the continuous upgrading of the thrust-to-weight ratio of aero-engines and the performance requirements of spacecraft, the new generation of high-temperature alloys faces dual technical challenges: firstly, they need to have long-term temperature resistance and structural stability in high-temperature service environments of 700℃~750℃; secondly, they need to have excellent hot working performance and formability to meet the precision manufacturing requirements of highly complex lightweight components, thereby breaking through the technical bottleneck of traditional materials in the synergistic optimization of temperature resistance and structural performance.

[0004] To improve the overall performance and reduce the cost of nickel-based superalloys, current methods mainly focus on adding trace alloying elements and controlling the composition of main alloying elements. The influence of trace elements on the performance of nickel-based superalloys is extremely complex. Under different service conditions (such as temperature gradients, stress environments, or corrosive atmospheres), their effects can even be diametrically opposed. The mechanisms and patterns of action of trace elements differ significantly across different alloy systems. These differences are reflected at the microscopic level, such as atomic diffusion and phase interface behavior, and further manifest in the control of macroscopic properties such as alloy strength and oxidation resistance.

[0005] In summary, providing a nickel-based superalloy that simultaneously meets the requirements of hot-end components of aero-engines for service environments of 700℃~750℃ and hot workability has become an urgent problem to be solved. Summary of the Invention

[0006] In view of the above, the present invention aims to provide a nickel-based superalloy and its preparation method, to solve at least one of the following technical problems: existing nickel-based superalloys are difficult to simultaneously meet the requirements of the hot end components of aero engines for service environment of 700℃~750℃ and hot workability.

[0007] The objective of this invention is mainly achieved through the following technical solutions:

[0008] This invention provides a nickel-based superalloy, the composition of which, by mass percentage, comprises: Mg: 0.01%~0.02%, Cu: 0.005%~0.1%, Fe: 3.8%~4.8%, Nb: 1%~1.3%, Cr: 15.2%~16.5%, Co: 9%~9.98%, Mo: 2.7%~3.4%, W: 2.5%~2.9%, Al: 2.1%~2.5%, Ti: 3.3%~3.7%, Zr: 0.025%~0.05%, C: 0.01%~0.014%, B: 0.005%~0.014%, and nickel: balance.

[0009] Furthermore, the composition of the nickel-based superalloy, by mass percentage, is (Ti+Nb) / Al: 1.95~2.1.

[0010] Furthermore, the composition of the nickel-based superalloy, by mass percentage, includes: Mg: 0.01%~0.02%, Cu: 0.005%~0.05%, Fe: 3.9%~4.4%, Nb: 1.2%~1.3%, Cr: 15.5%~16.2%, Co: 9.5%~9.98%, Mo: 3.0%~3.3%, W: 2.6%~2.8%, Al: 2.2%~2.5%, Ti: 3.4%~3.65%, Zr: 0.035%~0.049%, C: 0.01%~0.014%, B: 0.006%~0.013%, and nickel: balance.

[0011] The present invention also provides a method for preparing the above-mentioned nickel-based superalloy, comprising the following steps: Step 1: Prepare the raw materials according to the above alloy composition, and then perform vacuum induction melting, electroslag remelting and vacuum arc remelting in sequence to obtain alloy ingots. Then, perform homogenization annealing on the alloy ingots. Step 2: Hot deformation of the homogenized annealed alloy ingot is carried out at a temperature T. The deformation amount of each pass is not less than 30%. After each pass of deformation, the ingot is returned to the furnace and heated to T-20~40℃ and held for 2~4 hours until the alloy ingot is completely warmed up before the next pass of deformation is carried out. Step 3: Perform solution heat treatment and aging heat treatment on the hot-deformed alloy.

[0012] Furthermore, in step 1, the holding temperature for homogenization annealing is 1170℃~1190℃.

[0013] Furthermore, in step 2, T is 1000~1190℃.

[0014] Furthermore, in step 2, the deformation amount for each pass is 30% to 60%.

[0015] Furthermore, in step 2, the final cumulative true strain is 1.2~1.8.

[0016] Furthermore, in step 3, the specific process of solution treatment is as follows: after holding at 1000~1170℃ for 180~360min, it is cooled with oil.

[0017] Furthermore, in step 3, the specific process of aging treatment is as follows: first, keep it at 630~670℃ for 20~28h and then air cool it, and then keep it at 740~780℃ for 12~20h and then air cool it.

[0018] Compared with the prior art, the present invention can achieve at least one of the following beneficial effects: The nickel-based superalloy of this invention optimizes the hot working properties and long-term service stability of the alloy by adding trace amounts of Cu, and precisely controls the low Al+Ti content to enhance the strengthening phase γ. The content of the phase was reduced from 45% to 42%, thereby improving hot working properties; precise control of the W and Mo contents enhanced the contribution of the solid solution strengthening mechanism to performance, thus increasing the service temperature and γ-ray intensity. The stability of the phase at high temperatures was improved; the addition of Nb to the phase enhanced the γ-phase stability. The precipitation rate of the γ phase significantly improved the thermoplastic properties; strict control of the proportion and size of the reinforcing phase enabled the γ phase to be released more readily. The phase exhibits higher strength; the reduction of Co content and the addition of a large amount of Fe further reduce costs, while precise control of the composition of various trace elements improves the overall performance of the alloy. By precisely controlling the content and matching quantitative relationships of each element, it is possible to obtain a low-cost nickel-based superalloy that can meet the service environment of 700℃~750℃ and has good hot working performance.

[0019] In the preparation method of the nickel-based superalloy forging of the present invention, by precisely controlling the process parameters of each step, the microstructure of the nickel-based superalloy forging is ensured to mainly include equiaxed austenite grains and uniformly distributed carbides, as well as dispersed γ' phase, with a volume fraction of γ' phase of 40% or more, for example 40% to 43%, no η phase, and a grain size of 7.5 or more, wherein the equivalent diameter of the primary γ' phase is 0.9 μm to 2.1 μm, and the volume fraction of the primary γ' phase is 8% to 14%.

[0020] The nickel-based superalloy of this invention is low in cost and exhibits excellent high-temperature mechanical properties and hot working performance, meeting the requirements for aircraft engines operating at 700-750℃. For example, its properties are as follows: 700℃ performance: tensile strength ≥ 1330MPa (e.g., 1336-1360MPa); yield strength ≥ 1150MPa (e.g., 1160-1180MPa); elongation after fracture ≥ 12% (e.g., 12.5%-16%); reduction of area ψ ≥ 11% (e.g., 11%-14%); 750℃ performance: tensile strength ≥ 1210MPa (e.g., 1216-1250MPa); yield strength ≥ 1060MPa (e.g., 1069-1090MPa); elongation after fracture ≥ 9% (e.g., 9%-11%). The reduction of area ψ is ≥10% (e.g., 10% to 12%); performance at 750℃ / 450MPa: duration ≥135h, e.g., 136 to 302h; elongation after fracture ≥12%, e.g., 12% to 17%. The above nickel-based superalloys can be used in aero-engines.

[0021] The nickel-based superalloy of the present invention has a wide hot working temperature range, for example, the hot working temperature range reaches above 180°C, and it has good hot working performance.

[0022] Other features and advantages of the invention will be set forth in the description which follows, and will be apparent in part from the description, or may be learned by practicing the invention. The objects and other advantages of the invention may be realized and obtained by means of what is particularly pointed out in the written description and the accompanying drawings. Attached Figure Description

[0023] The accompanying drawings are for illustrative purposes only and are not intended to limit the invention. Throughout the drawings, the same reference numerals denote the same parts.

[0024] Figure 1 The grain characteristics of the nickel-based superalloy prepared in Example 1; Figure 2 It is the intragranular γ-ray of the nickel-based superalloy prepared in Example 1 Phase reinforcement phase morphology. Detailed Implementation

[0025] The preferred embodiments of the present invention will now be described in detail with reference to the accompanying drawings, which form part of the present invention and, together with the embodiments of the present invention, serve to illustrate the principles of the present invention.

[0026] This invention provides a nickel-based superalloy, wherein the components of the nickel-based superalloy, by mass percentage, include: Mg: 0.01%~0.02%, Cu: 0.005%~0.1%, Fe: 3.8%~4.8%, Nb: 1%~1.3%, Cr: 15.2%~16.5%, Co: 9%~9.98%, Mo: 2.7%~3.4%, W: 2.5%~2.9%, Al: 2.1%~2.5%, Ti: 3.3%~3.7%, Zr: 0.025%~0.05%, C: 0.01%~0.014%, B: 0.005%~0.014%, and nickel: balance.

[0027] Specifically, the composition of the above-mentioned nickel-based superalloy, by mass percentage, is (Ti+Nb) / Al: 1.95~2.1.

[0028] The following details the function and dosage selection of the components contained in this invention: Mg: Magnesium has a small atomic radius and typically preferentially agglomerates at grain boundaries. When combined with impurities such as sulfur and oxygen, it forms stable compounds (e.g., MgS, MgO), reducing the precipitation of brittle grain boundary phases such as sulfides and oxides, thereby enhancing grain boundary strength and improving the alloy's creep resistance and high-temperature plasticity. Trace amounts of magnesium can act as grain refiners, inhibiting abnormal grain growth, optimizing microstructure uniformity during hot working, and improving the alloy's fatigue performance and crack propagation resistance. In the composition system of this invention, magnesium microalloying effectively reduces the dynamic recrystallization activation energy, lowers hot deformation rheological stress, and significantly improves forging capability. Magnesium slows down the high-temperature oxidation rate by promoting the formation of dense oxide films (e.g., Cr2O3, Al2O3), especially in cyclic oxidation environments above 800°C. Magnesium can affect the size and distribution of the γ' phase (Ni3Al), delaying the coarsening process while inhibiting the formation of the Laves phase, thus improving long-term high-temperature stability. Further research revealed that in the system of this invention, a content of Mg of more than 0.02% can trigger the precipitation of harmful phases. Taking all factors into consideration, the content of magnesium in this invention is controlled between 0.01% and 0.02%.

[0029] In Cu-based superalloys, copper can partially dissolve in the nickel matrix, providing solid solution strengthening and improving the alloy's mid-temperature strength while maintaining a certain level of toughness. This optimizes the alloy's plastic deformation capacity, increases the hot working window, and reduces the tendency of alloy bars to crack during billet preparation, forging, and rolling. Cu can also adjust the precipitation behavior of the γ' strengthening phase, thereby optimizing long-term high-temperature service performance. Furthermore, the addition of copper significantly enhances the alloy's corrosion resistance in reducing acid environments (such as sulfuric acid and hydrochloric acid) and sulfur-containing media. Copper improves the material's stability in complex environments by promoting the formation of passivation films or inhibiting localized corrosion. Further research has revealed that in the system of this invention, a copper content exceeding 0.1% leads to decreased phase stability and induces interfacial embrittlement. Therefore, considering all factors, the copper content in this invention is controlled between 0.005% and 0.1%.

[0030] Fe: Iron is an inexpensive metallic element. Partially replacing expensive nickel and cobalt can significantly reduce alloy costs while maintaining a balance in high-temperature performance. Iron can dissolve in the nickel matrix, enhancing matrix strength, especially providing moderate solid solution strengthening in the 500℃~800℃ range. Iron improves the long-term high-temperature microstructure stability of the alloy by inhibiting the precipitation of harmful phases (such as the Laves phase) and optimizing the stability of the γ' phase (Ni3Al). Iron works synergistically with elements such as chromium and aluminum to promote the formation of dense Cr2O3 and Al2O3 oxide films, slowing down the high-temperature oxidation rate. Appropriate amounts of iron can regulate the alloy's plasticity and resistance to hot deformation, reducing the tendency to crack during hot rolling and forging. Excessive iron substitution for cobalt will reduce the overall performance of the alloy; therefore, considering all factors, the iron content in this invention is controlled between 3.8% and 4.8%.

[0031] Nb: Niobium combines with carbon to form stable MC-type carbides, which preferentially distribute at grain boundaries, inhibiting grain boundary migration and the precipitation of harmful phases, while reducing chromium carbide (Cr). 23 The coarsening of C6 enhances grain boundary stability. Niobium atoms partially dissolve in the nickel matrix, causing lattice distortion due to atomic size differences, thus improving matrix strength, especially maintaining good resistance to deformation at high temperatures. Niobium refines the weld microstructure, inhibits hot crack formation, and improves the alloy's weldability, making it suitable for complex components requiring weld forming. Niobium inhibits excessive coarsening of the γ' phase, delaying microstructural degradation during long-term high-temperature service. However, niobium content exceeding 1.3% in this alloy may lead to segregation and metallurgical defects, while also increasing alloy cost; therefore, the niobium content is controlled between 1% and 1.3%.

[0032] Cr: Chromium oxidizes at medium to high temperatures to form a dense Cr₂O₃ layer, which blocks oxygen diffusion inward and significantly reduces the oxidation rate. Chromium also inhibits sulfide corrosion, thus ensuring the long-term service life of alloy parts in marine saline-gas corrosive environments. Chromium's solubility in the nickel matrix is ​​approximately 30%, and it can improve medium- to high-temperature strength by inducing lattice distortion. In the system of this invention, a chromium content exceeding 16.5% will cause harmful phases to precipitate in the alloy; therefore, the chromium content is controlled between 15.2% and 16.5%.

[0033] Cobalt (Co) is one of the main matrix-forming elements in nickel-based superalloys. It enhances the creep resistance and endurance strength of the alloy during high-temperature service by reducing the stacking fault energy of the matrix. Furthermore, cobalt reduces the alloy's deformation resistance, improves forging plasticity during hot deformation, and reduces the cracking risk of highly alloyed materials. However, due to the high price of cobalt, its content needs to be controlled between 9% and 9.98%.

[0034] Mo: As a key solid solution strengthening element in nickel-based superalloys, molybdenum significantly improves the medium- and high-temperature strength of the matrix by severely distorting the nickel matrix lattice and hindering dislocation movement. Moreover, molybdenum has a lower density and cost compared to other heavy elements, making it a good element for high-temperature structural components that balance strength and cost. However, a molybdenum content of more than 3.4% in this alloy will lead to Laves phase segregation. Therefore, the molybdenum content is controlled between 2.7% and 3.4%.

[0035] W: The core role of tungsten in nickel-based superalloys comes from its ability to provide excellent performance in ultra-high temperature environments through extreme solid solution strengthening, improved creep resistance and optimized microstructure stability. However, the addition of tungsten will greatly increase the density and cost of the alloy, resulting in excessive weight and high cost of aerospace rotating parts. Therefore, the tungsten content is controlled between 2.5% and 2.9%.

[0036] Al: Aluminum is the core element for γ' phase strengthening and high-temperature oxidation resistance. Its role directly affects the high-temperature strength limit and life of the alloy. It provides medium and high temperature strength by inhibiting dislocation movement through γ' phase precipitation strengthening. However, the aluminum-titanium ratio must be strictly controlled to prevent η phase precipitation. Therefore, this invention controls the aluminum content to be between 2.1% and 2.5%.

[0037] Ti: Titanium is a key regulating element for strengthening the γ' phase and maintaining structural stability. It synergistically forms the Ni3(Al,Ti) ordered phase with aluminum, which significantly enhances the high-temperature strength of the alloy. At the same time, it optimizes the aluminum-titanium ratio and suppresses the η phase. Therefore, the titanium content is controlled between 3.3% and 3.7%.

[0038] (Ti+Nb) / Al: Aluminum, titanium, and niobium are typical strengthening phase-forming elements in nickel-based superalloys. Generally, the mechanical properties and microstructure are balanced by controlling the content of these three elements. The ratio of (Ti+Nb) / Al can significantly affect the mismatch between γ / γ', thereby affecting its high-temperature performance. The inventors have found through in-depth research that in the system of this invention, it is necessary to control (Ti+Nb) / Al > 1.95 to achieve higher mechanical properties. When (Ti+Nb) / Al ≥ 2.1, the η phase is very likely to precipitate in the alloy of this invention, which significantly reduces the alloy performance. Therefore, it is necessary to control (Ti+Nb) / Al between 1.95 and 2.1.

[0039] Zr: Zirconium is a key grain boundary strengthening trace element, generally enriched at grain boundaries, reducing sulfur concentration at grain boundaries and significantly improving the high-temperature creep life and process stability of the alloy. However, if the zirconium content in this alloy exceeds 0.04%, brittle ZrC points will form, which are prone to cracking and affect the alloy life. Therefore, the zirconium content is controlled between 0.025% and 0.05%.

[0040] C: Carbon is generally strengthened through carbides (MC / M) 23 C6) and grain boundary regulation enhance the high-temperature strength and microstructure stability, prevent grain boundary slip and crack propagation, improve the fatigue life of the alloy, and enhance its fatigue plasticity. Carbides can also act as heterogeneous nucleation sites, promoting grain refinement. However, excessive carbon content can lead to the formation of large primary carbides in the alloy, thereby reducing the fatigue life. Therefore, in this invention, the carbon content is controlled between 0.01% and 0.014%.

[0041] B: Boron works synergistically with carbon and zirconium to strengthen grain boundaries. Boron can significantly improve creep life, enrich grain boundaries, reduce grain boundary energy, inhibit grain boundary slip and crack initiation, and optimize microstructure stability. Considering all factors, the boron content is between 0.005% and 0.014%.

[0042] Specifically, in order to further improve the overall performance of the above-mentioned nickel-based superalloy, the composition of the above-mentioned nickel-based superalloy, by mass percentage, includes: Mg: 0.01%~0.02%, Cu: 0.005%~0.05%, Fe: 3.9%~4.4%, Nb: 1.2%~1.3%, Cr: 15.5%~16.2%, Co: 9.5%~9.98%, Mo: 3.0%~3.3%, W: 2.6%~2.8%, Al: 2.2%~2.5%, Ti: 3.4%~3.65%, Zr: 0.035%~0.049%, C: 0.01%~0.014%, B: 0.006%~0.013%, and nickel: balance.

[0043] This invention also provides a method for preparing the above-mentioned nickel-based superalloy, comprising the following steps: Step 1: Prepare the raw materials according to the above alloy composition, and then perform vacuum induction melting, electroslag remelting and vacuum arc remelting in sequence to obtain alloy ingots. Then, perform homogenization annealing on the alloy ingots. Step 2: Hot deformation of the homogenized annealed alloy ingot is carried out at T=1000~1190℃. The deformation amount of each pass is not less than 30%. After each deformation, the ingot is returned to the furnace and heated to T-20~40℃ and held for 2~4 hours until the alloy ingot is completely warmed up before the next deformation pass is carried out. The final cumulative true strain is 1.2~1.8. Step 3: Perform solution heat treatment and aging heat treatment on the hot-deformed alloy.

[0044] Specifically, in step 1 above, excessively high homogenization annealing temperatures can cause the alloy ingot to approach the initial melting temperature of the alloy, resulting in localized melting (overheating) and material scrap. Insufficient temperatures cannot effectively promote atomic diffusion, dissolve harmful phases such as the Laves phase and low-melting-point eutectic phase, or eliminate dendritic segregation in the ingot or billet. Insufficient time cannot guarantee that solute atoms such as Nb, Ti, and Mo can fully diffuse from the interdendritic space to the dendritic trunk, dissolve harmful phases, and homogenize the composition. Excessive time leads to excessively high costs. Therefore, the holding temperature for homogenization annealing is controlled at 1170℃~1190℃ (e.g., 1170℃, 1180℃, 1190℃); the homogenization time is 20~60h, e.g., 20h, 30h, 40h, 50h, 60h.

[0045] Specifically, in step 2 above, excessive deformation per pass can lead to uneven temperature distribution between the core and outer edge of the alloy ingot, potentially causing cracking. Insufficient deformation per pass, on the other hand, cannot provide enough energy to break down the original microstructure. Therefore, the deformation per pass should be controlled to be no less than 30%, for example, 30% to 60%, preferably 30% to 40%.

[0046] Specifically, in step 2 above, an excessively large final cumulative true strain will result in excessively fine grains, affecting the alloy's subsequent heat treatment processes and its creep performance during service. Conversely, an insufficient amount of deformation will prevent the alloy ingot from completely transforming from a cast structure to a uniform, fine-grained forged structure. Therefore, the final cumulative true strain should be controlled to be 1.2~1.8, for example, 1.2~1.6.

[0047] Specifically, in step 2 above, the alloy of the present invention has a wide hot working range, which is beneficial for selecting and controlling the processing temperature in industrial-scale production, reducing process sensitivity, and improving yield. For example, the hot working range of the alloy of the present invention is above 180°C, such as 180~200°C. When the hot working range is narrow, frequent reheating in the furnace is required during deformation, which can easily lead to abnormal growth of microstructure, increasing the risk of cracking, and resulting in low production efficiency, extremely long processing time, and high costs. Therefore, a wide hot working range is of great significance for the preparation of high-temperature alloys.

[0048] Specifically, in step 3 above, the solution treatment process is as follows: hold at 1000~1170℃ (e.g., 1000℃, 1020℃, 1050℃, 1070℃, 1100℃, 1120℃, 1150℃, 1170℃) for 180~360min (e.g., 180min, 200min, 220min, 250min, 280min, 300min, 330min, 350min, 360min) and then oil cool.

[0049] Specifically, in step 3 above, the aging process is as follows: first, keep it at 630~670℃ for 20~28 hours and then air cool it, and then keep it at 740~780℃ for 12~20 hours and then air cool it.

[0050] Specifically, the aforementioned nickel-based superalloys have a wide hot working temperature range, for example, reaching over 180°C, and exhibit good hot working performance.

[0051] Specifically, the microstructure of the aforementioned nickel-based superalloy mainly includes equiaxed austenite grains and uniformly distributed carbides, as well as a dispersed γ' phase. The volume fraction of the γ' phase is above 40%, for example, 40% to 43%, and there is no η phase. The grain size reaches grade 7.5 or above, for example, grade 7.5 to 8.5. The equivalent diameter of the primary γ' phase is 0.9 μm to 2.1 μm, and the volume fraction of the primary γ' phase is 8% to 14%.

[0052] Specifically, the aforementioned nickel-based superalloys can meet the material requirements for rotating components of aero-engines operating at 700~750℃, with the following properties: 700℃ properties: tensile strength ≥ 1330MPa (e.g., 1336~1360MPa); yield strength ≥ 1150MPa (e.g., 1160~1180MPa); elongation after fracture ≥ 12% (e.g., 12.5%~16%); reduction of area ψ ≥ 11% (e.g., 11%~14%); 750℃ properties: tensile strength ≥ 1210MPa (e.g., 1216~1250MPa); yield strength ≥ 1060MPa (e.g., 1069~1090MPa); elongation after fracture ≥ 9% (e.g., 9%~11%). Reduction of area ψ ≥ 10% (e.g., 10% to 12%); 750℃ / 450MPa performance: duration ≥ 135h, e.g., 136 to 302h; elongation after fracture ≥ 12%, e.g., 12% to 17%.

[0053] The present invention also provides applications of the above-mentioned nickel-based superalloy, which can be used in aero engines.

[0054] The nickel-based superalloy of this invention optimizes the hot working properties and long-term service stability of the alloy by adding trace amounts of Cu, and precisely controls the low Al+Ti content to enhance the strengthening phase γ. The content of the phase was reduced from 45% to 42%, thereby improving hot working properties; precise control of the W and Mo contents enhanced the contribution of the solid solution strengthening mechanism to performance, thus increasing the service temperature and γ-ray intensity. The stability of the phase at high temperatures was improved; the addition of Nb to the phase enhanced the γ-phase stability. The precipitation rate of the γ phase significantly improved the thermoplastic properties; strict control of the proportion and size of the reinforcing phase enabled the γ phase to be released more readily. The phase exhibits higher strength; the reduction of Co content and the addition of a large amount of Fe further reduce costs, while precise control of the composition of various trace elements improves the overall performance of the alloy. By precisely controlling the content and matching quantitative relationships of each element, it is possible to obtain a low-cost nickel-based superalloy that can meet the service environment of 700℃~750℃ and has good hot working performance.

[0055] In the preparation method of the nickel-based superalloy forging of the present invention, by precisely controlling the process parameters of each step, the microstructure of the nickel-based superalloy forging is ensured to mainly include equiaxed austenite grains and uniformly distributed carbides, as well as dispersed γ' phase, with a volume fraction of γ' phase of 40% or more, for example 40% to 43%, no η phase, and a grain size of 7.5 or more, wherein the equivalent diameter of the primary γ' phase is 0.9 μm to 2.1 μm, and the volume fraction of the primary γ' phase is 8% to 14%.

[0056] The nickel-based superalloy of this invention is low in cost and exhibits excellent high-temperature mechanical properties and hot working performance, meeting the requirements for aircraft engines operating at 700-750℃. For example, its properties are as follows: 700℃ performance: tensile strength ≥ 1330MPa (e.g., 1336-1360MPa); yield strength ≥ 1150MPa (e.g., 1160-1180MPa); elongation after fracture ≥ 12% (e.g., 12.5%-16%); reduction of area ψ ≥ 11% (e.g., 11%-14%); 750℃ performance: tensile strength ≥ 1210MPa (e.g., 1216-1250MPa); yield strength ≥ 1060MPa (e.g., 1069-1090MPa); elongation after fracture ≥ 9% (e.g., 9%-11%). The reduction of area ψ is ≥10% (e.g., 10% to 12%); performance at 750℃ / 450MPa: duration ≥135h, e.g., 136 to 302h; elongation after fracture ≥12%, e.g., 12% to 17%. The above nickel-based superalloys can be used in aero-engines.

[0057] The nickel-based superalloy of the present invention has a wide hot working temperature range, for example, the hot working temperature range reaches above 180°C, and it has good hot working performance.

[0058] The advantages of precise control of the composition and process parameters of the high-temperature alloy of the present invention will be demonstrated below with specific embodiments and comparative examples.

[0059] Examples 1-3 of the present invention provide a nickel-based superalloy and its preparation method. The chemical composition of the nickel-based superalloys of Examples 1-3 is shown in Table 1.

[0060] The methods for preparing steel in Examples 1-3 include: Step 1: Prepare the raw materials according to the above alloy composition, and then perform vacuum induction melting, electroslag remelting and vacuum arc remelting in sequence to obtain alloy ingots. Then, perform homogenization annealing on the alloy ingots. The holding temperature for homogenization annealing is 1170℃~1190℃ and the homogenization time is 20~60h. Step 2: Hot deformation of the homogenized annealed alloy ingot is carried out at 1000~1190℃, with a deformation amount of not less than 30% per pass. After each deformation, the ingot is returned to the furnace and heated to 1000~1190℃ and held for 2~4 hours until the alloy ingot is completely warmed up before the next deformation pass is carried out. The final cumulative true strain is 1.2~1.8. Step 3: Perform solution heat treatment and aging heat treatment on the hot-deformed alloy. The specific process of solution treatment is: hold at 1000~1170℃ for 180~360min and then oil cool. The specific process of aging treatment is: first hold at 630~670℃ for 20~28h and then air cool, and then hold at 740~780℃ for 12~20h and then air cool.

[0061] The specific process parameters of Examples 1-3 are shown in Table 2; the grain size and precipitated phase of Examples 1-3 are shown in Table 3; the performance test results of Examples 1-3 at 700℃ and 750℃ are shown in Tables 4 and 5; and the hot working range of Examples 1-3 is shown in Table 6.

[0062] Figure 1 The grain characteristics of the nickel-based superalloy prepared in Example 1; Figure 2 It is the intragranular γ-ray of the nickel-based superalloy prepared in Example 1 Phase reinforcement phase morphology.

[0063] The microstructure of Examples 1-3 mainly includes equiaxed austenite grains and uniformly distributed carbides, as well as a dispersed γ' phase. The volume fraction of the γ' phase is above 40%, for example, 40% to 43%, and there is no η phase. The grain size reaches grade 7.5 or above, wherein the equivalent diameter of the primary γ' phase is 0.9 μm to 2.1 μm, and the volume fraction of the primary γ' phase is 8% to 14%.

[0064] The inventors conducted extensive research during the research process, and some suboptimal solutions are presented here as comparative examples.

[0065] Comparative Example 1 This comparative example provides a nickel-based superalloy and its preparation method. The composition is shown in Table 1. The preparation method is the same as that in Example 1, and will not be repeated here.

[0066] Comparative Example 2 This comparative example provides a nickel-based superalloy and its preparation method. Its composition is shown in Table 1 above. The preparation method is the same as that in Example 1, and will not be repeated here.

[0067] Comparative Example 3 This comparative example provides a nickel-based superalloy and its preparation method. Its composition is shown in Table 1 above. The preparation method is the same as that in Example 1, and will not be repeated here.

[0068] Comparative Example 4 This comparative example provides a nickel-based superalloy and its preparation method. Its composition is shown in Table 1 above. In the preparation method, the deformation amount per pass is about 15%, and the cumulative true strain is 0.6.

[0069] The microstructure and main properties of the comparative steel are shown in Tables 3-6 below.

[0070] As can be seen from the results in Tables 3-7, the embodiments of the present invention precisely control the process parameters of the preparation method by accurately determining the types and contents of elements in the components and the synergistic quantitative relationships, thus ensuring that the nickel-based high-temperature alloy can meet the material requirements for rotating parts of aero-engines in service at 700~750℃.

[0071] Table 1 Chemical composition (mass %) of the examples and comparative examples

[0072] Table 2 Specific process parameters

[0073] Table 3. Grain size and precipitates of the examples and comparative examples.

[0074] Table 4. Properties of the alloy at 700℃

[0075] Table 5 Properties of the alloy at 750℃

[0076] Table 6. Properties of the alloy at 750℃ / 450MPa

[0077] Table 7 Hot working temperature range

[0078] The above description is only a preferred embodiment of the present invention, but the scope of protection of the present invention is not limited thereto. Any changes or substitutions that can be easily conceived by those skilled in the art within the scope of the technology disclosed in the present invention should be included within the scope of protection of the present invention.

Claims

1. A nickel-based superalloy, characterized in that, The composition of the nickel-based superalloy, by mass percentage, includes: Mg: 0.01%~0.02%, Cu: 0.005%~0.1%, Fe: 3.8%~4.8%, Nb: 1%~1.3%, Cr: 15.2%~16.5%, Co: 9%~9.98%, Mo: 2.7%~3.4%, W: 2.5%~2.9%, Al: 2.1%~2.5%, Ti: 3.3%~3.7%, Zr: 0.025%~0.05%, C: 0.01%~0.014%, B: 0.005%~0.014%, and nickel: balance.

2. The nickel-based superalloy according to claim 1, characterized in that, The composition of the nickel-based superalloy, by mass percentage, is (Ti+Nb) / Al: 1.95~2.

1.

3. The nickel-based superalloy according to claim 1, characterized in that, The composition of the nickel-based superalloy, by mass percentage, includes: Mg: 0.01%~0.02%, Cu: 0.005%~0.05%, Fe: 3.9%~4.4%, Nb: 1.2%~1.3%, Cr: 15.5%~16.2%, Co: 9.5%~9.98%, Mo: 3.0%~3.3%, W: 2.6%~2.8%, Al: 2.2%~2.5%, Ti: 3.4%~3.65%, Zr: 0.035%~0.049%, C: 0.01%~0.014%, B: 0.006%~0.013%, and nickel: balance.

4. A method for preparing a nickel-based superalloy according to any one of claims 1 to 3, characterized in that, Includes the following steps: Step 1: Prepare the raw materials according to the above alloy composition, and then perform vacuum induction melting, electroslag remelting and vacuum arc remelting in sequence to obtain alloy ingots. Then, perform homogenization annealing on the alloy ingots. Step 2: Hot deformation of the homogenized annealed alloy ingot is carried out at a temperature T. The deformation amount of each pass is not less than 30%. After each pass of deformation, the ingot is returned to the furnace and heated to T-20~40℃ and held for 2~4 hours until the alloy ingot is completely warmed up before the next pass of deformation is carried out. Step 3: Perform solution heat treatment and aging heat treatment on the hot-deformed alloy.

5. The preparation method according to claim 4, characterized in that, In step 1, the holding temperature for homogenization annealing is 1170℃~1190℃.

6. The preparation method according to claim 4, characterized in that, In step 2, T is 1000~1190℃.

7. The preparation method according to claim 4, characterized in that, In step 2, the deformation amount for each pass is 30% to 60%.

8. The preparation method according to claim 4, characterized in that, In step 2, the final cumulative true strain is 1.2~1.

8.

9. The preparation method according to any one of claims 4 to 8, characterized in that, In step 3, the specific process of solution treatment is as follows: after holding at 1000~1170℃ for 180~360min, it is cooled with oil.

10. The preparation method according to any one of claims 4 to 8, characterized in that, In step 3, the specific process of aging treatment is as follows: first, keep it at 630~670℃ for 20~28h and then air cool it, and then keep it at 740~780℃ for 12~20h and then air cool it.

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