Nickel-based high-temperature alloy and preparation method thereof

Abstract

The invention discloses a nickel-based superalloy and a preparation method thereof, and relates to the technical field of nickel-based superalloys.The nickel-based superalloy comprises a basic alloy and a coating alloy; the coating alloy is molten brazing filler metal; the brazing filler metal is coated on the surface of the basic alloy in a molten state to form the nickel-based superalloy. According to the invention, through the collaborative design of the basic alloy and the cladding alloy, the synergistic interaction of various structural characteristics is realized; wherein Nb and Ta elements in the basic alloy and B and Rh elements in the cladding alloy form a multi-gradient strengthening structure, the structure shows the dynamic balance characteristic in the high-temperature service process, and a formed niobium diboride hard phase network provides mechanical strengthening and further forms an element diffusion barrier layer; and adverse migration of matrix elements to the surface at high temperature is effectively prevented.

Description

Technical Field

[0001] The present invention relates to the technical field of nickel-based superalloys, and particularly to a nickel-based superalloy and a preparation method thereof. Background Art

[0002] Nickel-based superalloys are indispensable key materials in high-temperature service environments such as aeroengines, industrial gas turbines, and nuclear energy equipment. Among them, nickel-based superalloys have a wide range of applications in the medium-temperature application field due to their good cost-effectiveness ratio and workability.

[0003] However, there are several significant deficiencies in the existing nickel-based superalloy technology. First, the interfacial bonding strength formed by the traditional coating process is insufficient, resulting in easy spalling failure of the coating under high-temperature cyclic conditions; second, it is difficult to control the uniform distribution of heavy elements such as Nb and Ta by the conventional melting process, causing large performance fluctuations; third, the strength retention rate of traditional superalloys in high-temperature environments is relatively low, generally only able to maintain 40-50% of the room-temperature strength, severely restricting their high-end applications. Summary of the Invention

[0004] In view of the problems existing in the existing nickel-based superalloys and their preparation methods, the present invention is proposed.

[0005] To solve the above technical problems, the present invention provides the following technical solutions:

[0006] In a first aspect, an embodiment of the present invention provides a nickel-based superalloy, including a base alloy and a cladding alloy; the cladding alloy is a molten solder; the solder is coated on the surface of the base alloy in a molten state to form the nickel-based superalloy.

[0007] The present invention adopts a composite structure design of a base alloy and a molten solder cladding, realizing hierarchical optimization of material functions. This structure enables the base alloy to bear the main mechanical load, while the cladding alloy provides surface protection, solving the technical contradiction that it is difficult for a single material to simultaneously meet high-temperature strength and surface oxidation resistance.

[0008] As a preferred embodiment of the nickel-based superalloy of the present invention, wherein: the base alloy has the following components by mass percentage,

[0009] C ≤ 0.06%; Cr: 16.6 - 17.7%; Nb: 6.1 - 6.8%; N: 6.5 - 7.0%; Ta: 4.3 - 4.8%; Co ≤ 0.7%; Mn: 2.7 - 3.1%; Al: 1.0 - 1.3%; the balance is Fe.

[0010] As a preferred embodiment of the nickel-based superalloy of the present invention, wherein: the cladding alloy has the following components by mass percentage,

[0011] Ni: 35 - 45%; Cr: 10 - 15%; Fe: 15 - 20%; W: 3 - 5%; Ti: 1.0 - 1.5%; Rh: 15 - 20%; B: 0.5 - 2.0%; the balance is Zr.

[0012] As a preferred embodiment of the nickel - based superalloy of the present invention, wherein: element Nb in the base alloy and element B in the cladding alloy obtain a niobium diboride hard phase at the interface; the hard phase is distributed in a network pattern in the interface region and has a thickness of 5 - 20 μm.

[0013] The network - distributed niobium diboride hard phase obtained at the interface creates a unique gradient strengthening structure. This network distribution with a thickness of 5 - 20 μm enables the interface to have both high strength and sufficient toughness, solving the technical dilemma of the traditional brazed interface being "strong but brittle, tough but weak"; more importantly, this network - distributed hard phase exhibits a "self - anchoring" effect at high temperatures. As the temperature rises, the interface bonding strength increases instead of decreasing.

[0014] In a second aspect, the embodiments of the present invention provide a preparation method of a nickel - based superalloy, which includes the following steps:

[0015] Melting the base alloy in an inert atmosphere at a melting temperature of 1250 - 1450 °C for a melting duration of 0.5 - 1.5 h to obtain an alloy ingot blank;

[0016] Keeping the alloy ingot blank obtained by melting at a temperature of 1140 - 1150 °C for 0.5 - 1.5 h, and pressing it into a plate by a press. The thickness of the pressed plate is 1.2 - 1.5 mm;

[0017] Cooling the compressed plate to 650 °C at a cooling rate of 3 - 5 °C / min, and then cooling it to 200 - 300 °C at a cooling rate of 1 - 2 °C / min to obtain an alloy plate;

[0018] After coating the surface of the alloy plate with the cladding alloy, the nickel - based superalloy is obtained.

[0019] As a preferred embodiment of the preparation method of the nickel - based superalloy of the present invention, wherein: the preparation method of the base alloy includes:

[0020] Raw materials in terms of mass percentage: C ≤ 0.06%, Cr: 16.6 - 17.7%, Nb: 6.1 - 6.8%, N: 6.5 - 7.0%, Ta: 4.3 - 4.8%, Co ≤ 0.7%, Mn: 2.7 - 3.1%, Al: 1.0 - 1.3%, and the balance is Fe;

[0021] The raw materials are divided into two groups. The first group of raw materials includes Fe, C, and Cr, and the second group contains Nb, N, Ta, Co, Mn, and Al;

[0022] In a vacuum induction melting furnace, the first group of raw materials is heated to 1400 - 1450 °C and held for 20 - 30 min to obtain molten material a;

[0023] Meanwhile, in another vacuum induction melting furnace, the second group of raw materials is heated to 1400 - 1450 °C and held for 20 - 30 min to obtain molten material b;

[0024] Under the protection of an inert atmosphere, molten material a is poured into molten material b, stirred evenly and held for 10 - 20 min to obtain an alloy melt;

[0025] The uniformly mixed alloy melt is poured into a water-cooled steel mold lined with a ceramic cylinder at a speed of 200 - 300 mm / min, and left standing for 30 - 60 min to obtain the base alloy;

[0026] Among them, the material of the inner ceramic cylinder is alumina ceramic, and the inner wall thickness is 5 - 10 mm;

[0027] The water-cooled steel mold is cooled by circulating water.

[0028] As a preferred embodiment of the preparation method of the nickel-based superalloy of the present invention, wherein: the preparation method of the cladding alloy includes,

[0029] Raw materials of Ni: 35 - 45%, Cr: 10 - 15%, Fe: 15 - 20%, W: 3 - 5%, Ti: 1.0 - 1.5%, Rh: 15 - 20%, B: 0.5 - 2.0%, and the balance Zr are measured by mass percentage;

[0030] In an electron beam melting furnace, the raw materials are heated to 1400 - 1500 °C and held for 0.5 - 1.5 h under the protection of pure argon to obtain a molten cladding alloy;

[0031] The molten cladding alloy is cooled to 800 - 900 °C and kept in a molten state;

[0032] Using a vacuum magnetron sputtering coating equipment, under the condition that the sputtering power density is 10 - 20 W / cm 2 the molten cladding alloy is evenly coated on the surface of the base alloy plate at a sputtering speed of 150 - 200 s / min;

[0033] After coating, it is held in a vacuum environment at 1000 - 1100 °C for 20 - 30 s to complete the preliminary bonding of the cladding alloy and the base alloy.

[0034] As a preferred embodiment of the preparation method of the nickel-based superalloy of the present invention, wherein: after the preliminary bonding of the clad alloy and the base alloy, the following steps are further included.

[0035] The preliminarily bonded clad alloy and base alloy are heated at 2100 - 2400 °C for 2 - 3 h, and then air-cooled to 25 °C.

[0036] Then it is heated to 1050 °C and vacuum annealed for 40 - 80 min, and then air-cooled to 25 °C to obtain the nickel-based superalloy.

[0037] The beneficial effects of the present invention are as follows:

[0038] Through the collaborative design of the base alloy and the clad alloy, the synergistic effect of multiple structural features is achieved. Among them, the Nb and Ta elements in the base alloy and the B and Rh elements in the clad alloy form a multi-gradient strengthening structure, which exhibits a "dynamic balance" characteristic during high-temperature service. The formed niobium boride hard phase network not only provides mechanical strengthening but also forms an element diffusion barrier layer, effectively preventing the adverse migration of matrix elements to the surface at high temperatures; secondly, through a multi-stage cooling process (cooling at 3 - 5 °C / min to 650 °C, and then cooling at 1 - 2 °C / min to 200 - 300 °C) and plate pressing (1.2 - 1.5 mm), a directionally arranged Nb-rich nanoclusters are formed, and these clusters become the preferred sites for the growth of niobium boride during subsequent heat treatment, making the finally formed hard phase network present a three-dimensional interconnected structure instead of a random distribution in the traditional process, and the three-dimensional interconnected structure realizes the improvement of structural strength.

[0039] The high-content (15 - 20%) Rh element in the clad alloy forms a micro-scale interlocking structure with niobium boride at the interface. When microcracks appear in the hard phase network, the Rh element promotes the rapid diffusion of B, realizing the "self-healing" function; extending the service life of the material under high-temperature cyclic conditions, and the use of N element forms nano-scale nitrides in the base alloy, and these nitrides pin dislocations and refine grains, providing a microscopic basis for the niobium boride hard phase network and improving the thermal stability of the alloy.

[0040] Fe, C, Cr and Nb, N, Ta, Co, Mn, Al are melted separately in two groups and then mixed under the protection of an inert atmosphere, solving the problem that heavy elements such as Nb and Ta are prone to segregation at the bottom of the melt in the traditional process, resulting in non-uniform structure. The grouped melting of the present invention makes these elements in an active state before mixing, and form a uniform suspension after mixing, with a significant reduction in the segregation index; and the temperature gradient and composition gradient generated during the mixing of the two groups of melts induce the formation of a unique "core-shell structure" strengthening phase, which evolves into a high-density dislocation network during heat treatment, greatly improving the high-temperature strength and creep resistance of the material. Detailed implementation manners

[0041] To make the above objects, features and advantages of the present invention more obvious and understandable, the following provides a detailed description of the specific implementation manners of the present invention. Obviously, the described embodiments are part of the embodiments of the present invention, rather than all embodiments. All other embodiments obtained by those of ordinary skill in the art based on the embodiments of the present invention without creative efforts shall fall within the protection scope of the present invention.

[0042] In the following description, many specific details are set forth to fully understand the present invention. However, the present invention can also be implemented in other ways different from those described herein. Those skilled in the art can make similar extensions without departing from the connotation of the present invention. Therefore, the present invention is not limited by the specific embodiments disclosed below.

[0043] Secondly, the so-called "one embodiment" or "embodiment" herein refers to a specific feature, structure or characteristic that can be included in at least one implementation manner of the present invention. The phrase "in one embodiment" appearing in different places in this specification does not necessarily refer to the same embodiment, nor is it a separate or alternative embodiment that excludes other embodiments.

[0044] Embodiment 1

[0045] This embodiment provides a nickel-based superalloy, including a base alloy and a cladding alloy; the cladding alloy is a molten solder; the solder is coated on the surface of the base alloy in a molten state to form the nickel-based superalloy.

[0046] The base alloy has the following components by mass percentage:

[0047] C: 0.02%; Cr: 16.6%; Nb: 6.8%; N: 7.0%; Ta: 4.3%; Co: 0.2%; Mn: 2.7%; Al: 1.3%; the balance is Fe.

[0048] The cladding alloy has the following components by mass percentage:

[0049] Ni: 45%; Cr: 15%; Fe: 20%; W: 3%; Ti: 1.0%; Rh: 15%; B: 0.5%; the balance is Zr.

[0050] Element Nb in the base alloy and element B in the cladding alloy obtain a niobium diboride hard phase at the interface; the hard phase is distributed in a network pattern in the interface region and has a thickness of 5 μm.

[0051] Embodiment 2

[0052] This embodiment provides a nickel-based superalloy, including a base alloy and a clad alloy; the clad alloy is a molten solder; the solder is coated on the surface of the base alloy in a molten state to form the nickel-based superalloy.

[0053] The components of the base alloy by mass percentage are as follows:

[0054] C: 0.06%; Cr: 17.7%; Nb: 6.1%; N: 6.5%; Ta: 4.8%; Co: 0.7%; Mn: 3.1%; Al: 1.0%; the balance is Fe.

[0055] The components of the clad alloy by mass percentage are as follows:

[0056] Ni: 35%; Cr: 10%; Fe: 15%; W: 5%; Ti: 1.5%; Rh: 20%; B: 2%; the balance is Zr.

[0057] Element Nb in the base alloy and element B in the clad alloy obtain niobium diboride hard phase at the interface; the hard phase is distributed in a network pattern in the interface region and has a thickness of 10 μm.

[0058] Example 3

[0059] This embodiment provides a nickel-based superalloy, including a base alloy and a clad alloy; the clad alloy is a molten solder; the solder is coated on the surface of the base alloy in a molten state to form the nickel-based superalloy.

[0060] The components of the base alloy by mass percentage are as follows:

[0061] C: 0.04%; Cr: 17%; Nb: 6.4%; N: 6.8%; Ta: 4.5%; Co: 0.4%; Mn: 2.9%; Al: 1.2%; the balance is Fe.

[0062] The components of the clad alloy by mass percentage are as follows:

[0063] Ni: 40%; Cr: 13%; Fe: 18%; W: 4%; Ti: 1.3%; Rh: 18%; B: 1%; the balance is Zr.

[0064] Element Nb in the base alloy and element B in the clad alloy obtain niobium diboride hard phase at the interface; the hard phase is distributed in a network pattern in the interface region and has a thickness of 20 μm.

[0065] Comparative Example 1: Reduce the mass of Nb in the base alloy and B in the clad alloy. Specifically:

[0066] A nickel-based superalloy comprises a base alloy and a cladding alloy; the components of the base alloy by mass percentage are as follows,

[0067] C: 0.04%; Cr: 17%; Nb: 2.5%; N: 6.8%; Ta: 4.5%; Co: 0.4%; Mn: 2.9%; Al: 1.2%; the balance is Fe.

[0068] The components of the cladding alloy by mass percentage are as follows,

[0069] Ni: 40%; Cr: 13%; Fe: 18%; W: 4%; Ti: 1.3%; Rh: 18%; B: 0.1%; the balance is Zr.

[0070] Other components are the same as those in Example 3.

[0071] Comparative Example 2: There is no Rh in the cladding alloy, specifically expressed as:

[0072] A nickel-based superalloy comprises a base alloy and a cladding alloy; the components of the base alloy by mass percentage are as follows,

[0073] C: 0.04%; Cr: 17%; Nb: 6.4%; N: 6.8%; Ta: 4.5%; Co: 0.4%; Mn: 2.9%; Al: 1.2%; the balance is Fe.

[0074] The components of the cladding alloy by mass percentage are as follows,

[0075] Ni: 58%; Cr: 13%; Fe: 18%; W: 4%; Ti: 1.3%; Rh: 0%; B: 1%; the balance is Zr.

[0076] Other components are the same as those in Example 3.

[0077] Through the analysis of scanning electron microscopy (SEM) combined with energy dispersive spectroscopy (EDS), there are significant differences in the interfacial microstructures of the three samples:

[0078] Example 3: A complete and continuous niobium diboride hard phase network is formed at the interface, showing a typical three-dimensional interconnected distribution, and the network thickness is uniformly up to 20 μm. The hard phase particles are fine (100 - 150 nm) and uniformly distributed, and the lattice structure shows a hexagonal crystal system, forming a semi-coherent relationship with the matrix, the network coverage rate is above 95%, and the porosity is below 2%.

[0079] Comparative Example 1: Since the Nb content in the base alloy is low (2.5%) and the B content in the cladding is only 0.1%, only sporadic hard phase particles are formed at the interface, with uneven thickness (2 - 8 μm). The hard phases exist in the form of coarse particles (1 - 3 μm), isolated from each other, and cannot form a continuous network. The coverage rate is only 25 - 30%, with a large number of blank areas.

[0080] Comparative Example 2: Although the Nb content in the base alloy and the B content in the cladding are the same as in Example 3, lacking the catalytic effect of the Rh element, the niobium diboride phase at the interface is distributed in discontinuous flakes, with a thickness of 10 - 15 μm and large fluctuations. The sizes of the hard phases are uneven (200 nm - 2 μm), coarsening and agglomeration occur in some areas, the network continuity is poor, and there are obvious broken network areas.

[0081] The micro - tensile test technique was used to measure the interfacial shear strength at different temperatures, and the results are shown in Table 1;

[0082] Table 1: Shear strength parameter table at different temperatures

[0083] Test conditions Example 3 Comparative Example 1 Comparative Example 2 25℃ 210 MPa 55 MPa 105 MPa 600℃ 225 MPa 48 MPa 92 MPa 800℃ 240 MPa 35 MPa 78 MPa 1000℃ 255 MPa 22 MPa 65 MPa

[0084] It can be seen from Table 1 that the interfacial strength of Example 3 increases with the increase of temperature (positive temperature coefficient), which is due to the "self - anchoring" effect of the niobium diboride network at high temperatures; the strength of Comparative Example 1 is the lowest, only 40% of that at room temperature at 1000 °C, and serious interfacial softening occurs; although the performance of Comparative Example 2 is improved compared with Comparative Example 1, the high - temperature performance still decays significantly.

[0085] Isothermal oxidation tests (for 100 hours) were carried out in static air at 1100 °C for high - temperature oxidation behavior analysis, and the results are shown in Table 2;

[0086] Table 2: High - temperature oxidation parameter table

[0087] Test indicators Example 3 Comparative Example 1 Comparative Example 2 <![CDATA[Oxidation weight gain (mg / cm 2 )]]> 0.75 3.85 5.20 <![CDATA[Oxidation rate constant (mg 2 / cm 4 ·h)]]> 0.0056 0.148 0.270 Oxide layer structure Continuous and dense Multi-layered and loose Inner and outer double layer Spalling condition No spalling 15% area spalling Severe spalling at the edge

[0088] In Example 3, a dense (Rh,Cr)2O3 protective film with a thickness of only 8 - 10 μm was formed. The presence of the Rh element significantly reduced the growth rate of the oxide film and improved the adhesion. When micro - cracks appeared on the surface, the B element quickly diffused to form B2O3 to fill the defects, showing excellent self - healing ability; although Comparative Example 1 contains the Rh element, the insufficient B content results in the inability to form an effective interfacial barrier layer, and oxygen diffuses rapidly along the grain boundaries, forming a porous multi - layer oxide; Comparative Example 2 has no Rh element, and the formed Cr2O3 oxide film has poor bonding with the substrate, is prone to spalling under thermal cycling conditions, and loses its protective effect.

[0089] The thermal cycle test was carried out between 300 - 1050 °C. Each cycle included heating for 15 minutes, holding for 20 minutes, and cooling for 15 minutes. The results are shown in Table 3;

[0090] Table 3: Thermal cycle stability performance parameters

[0091] Performance indicators Example 3 Comparative Example 1 Comparative Example 2 First crack appearance (number of cycles) >500 38 125 Interface integrity (after 250 cycles) 100% intact Completely separated 60% peeling Dimensional stability (deformation rate) <0.5% 3.8% 2.2% Interface strength retention rate (after 500 cycles) 96% - -

[0092] Example 3 provides excellent stress buffering ability through a complete niobium boride network. The Rh element improves the thermal expansion matching between the oxide film and the matrix, and the uniform 20 - μm - thick interface layer avoids stress concentration; Comparative Example 1 and Comparative Example 2 failed rapidly under thermal cycle conditions due to interface structure defects and poor oxide film quality.

[0093] Through the comparative analysis of Example 3 with Comparative Example 1 and Comparative Example 2, it is proved that the present invention has successfully achieved an ideal reticular structure of niobium boride hard phase by precisely controlling the contents of Nb, B, and Rh, and has reached an unprecedented level in key performance indicators such as interface strength, oxidation resistance, and thermal stability.

[0094] Example 4

[0095] This example provides a preparation method of a nickel - based superalloy, including the following steps A1 - A4,

[0096] A1: Melting the base alloy in an inert atmosphere at a melting temperature of 1300 °C for a melting duration of 1 h to obtain an alloy ingot blank;

[0097] A2: Holding the alloy ingot blank obtained by melting at a temperature of 1140 °C for 1 h, and pressing it into a plate with a press. The thickness of the pressed plate is 1.5 mm;

[0098] A3: Cooling the compressed plate at a cooling rate of 5 °C / min to 650 °C, and then cooling it at a cooling rate of 2 °C / min to 250 °C to obtain an alloy plate;

[0099] A4: After coating the surface of the alloy plate with a cladding alloy, a nickel - based superalloy is obtained.

[0100] The preparation method of the base alloy includes A1.1 - A1.6:

[0101] A1.1: Weighing each raw material of C, C, Nb, N, Ta, Co, Mn, Al, and the balance being Fe according to mass percentage;

[0102] A1.2: Dividing the raw materials into two groups. The first group of raw materials includes Fe, C, and Cr, and the second group includes Nb, N, Ta, Co, Mn, and Al;

[0103] A1.3: In a vacuum induction melting furnace, heat the first group of raw materials to 1400 °C and hold for 25 min to obtain molten material a;

[0104] A1.4: Meanwhile, in another vacuum induction melting furnace, heat the second group of raw materials to 1450 °C and hold for 25 min to obtain molten material b;

[0105] A1.5: Under the protection of an inert protective atmosphere, pour molten material a into molten material b, stir evenly and hold for 10 - 20 min to obtain an alloy melt;

[0106] A1.6: Pour the uniformly mixed alloy melt into a water-cooled steel mold lined with a ceramic cylinder at a speed of 250 mm / min, and let it stand for 0 min to obtain a base alloy;

[0107] Among them, the material of the inner ceramic cylinder is alumina ceramic, and the inner wall thickness is 5 mm; the water-cooled steel mold is cooled by circulating water.

[0108] The preparation method of the clad alloy includes B1 - B5:

[0109] B1: Measure each raw material of Ni, Cr, Fe, W, Ti, Rh, B, and the balance Zr by mass percentage;

[0110] B2: In an electron beam melting furnace, heat each raw material to 1500 °C and hold for 1 h under the protection of pure argon to obtain a molten clad alloy;

[0111] B3: Cool the molten clad alloy to 900 °C and keep it in a molten state;

[0112] B4: Use a vacuum magnetron sputtering coating equipment. Under the condition that the sputtering power density is 10 W / cm 2 coat the molten clad alloy evenly on the surface of the base alloy plate at a sputtering speed of 150 s / min;

[0113] B5: After coating, keep it in a vacuum environment at 1000 °C for 30 s to complete the preliminary bonding of the clad alloy and the base alloy.

[0114] After the preliminary bonding of the clad alloy and the base alloy, the following steps C1 - C2 are also included:

[0115] C1: Heat the preliminarily bonded clad alloy and the base alloy at 2100 °C for 2 h, and then air-cool to 25 °C;

[0116] C2: Then heat up to 1050 °C and perform vacuum annealing for 40 min, and then air-cool to 25 °C to obtain a nickel-based superalloy.

[0117] Comparative Example 3: The base alloy was prepared by a traditional one-step melting process. All raw materials were added to a vacuum induction melting furnace for melting simultaneously without grouped melting. The total raw materials were directly held at 1400 °C for 30 min. A conventional casting mold was used without a special inner ceramic cylinder. The remaining process steps were the same as those in Example 4.

[0118] Comparative Example 4: Prepared by a rapid cooling process. Using the same grouped melting process as in Example 4, the alloy ingot blank obtained by melting was held at 1145 °C for 1 h and then pressed into a plate with a thickness of 1.3 mm. A conventional rapid cooling process was used to directly cool to room temperature at a rate of 25 °C / min without using a multi-stage precise cooling process. The remaining process steps were the same as those in Example 4.

[0119] Comparative Example 5: A nickel-based superalloy without ultra-high temperature post-treatment. Using the same grouped melting process as in Example 4 and the same multi-stage precise cooling process as in Example 4, after coating a cladding alloy on the surface of the base alloy plate by vacuum magnetron sputtering, it was held in a vacuum environment at 1100 °C for 30 s without undergoing ultra-high temperature treatment at 2100 °C, and only annealed at 1000 °C for 60 min. The remaining process steps were the same as those in Example 4.

[0120] It can be seen from the comparison of Comparative Examples 3 - 5 with Example 4:

[0121] Comparative Example 3 adopted a traditional one-step melting process, resulting in severe segregation of heavy elements and the inability to form a uniform network of interfacial hard phases. While in Example 4, the grouped melting process ensured the uniform distribution of heavy elements such as Nb and Ta, laying a microscopic foundation for subsequent interfacial reactions. This demonstrated the decisive role of grouped melting in the formation of high-performance interfacial structures.

[0122] Comparative Example 4 used a conventional rapid cooling, failing to form a Nb-rich surface region, resulting in uneven sizes and discontinuous distributions of interfacial hard phases. The multi-stage precise cooling process in Example 4 (cooling to 650 °C at 5 °C / min and then to 250 °C at 2 °C / min) promoted the directional migration of Nb elements to form surface enrichment, while significantly reducing residual stress, enabling the final material to exhibit excellent stability under thermal cycling conditions.

[0123] Although Comparative Example 5 adopted grouped melting and multi-stage cooling, the lack of ultra-high temperature post-treatment led to insufficient development of niobium boride phases at the interface and poor bonding with the matrix. The 2100 °C ultra-high temperature treatment in step C1 of Example 4 triggered controllable recrystallization in the interfacial region, enabling the niobium boride hard phase to form a semi-coherent relationship with the matrix, significantly reducing the interfacial energy, and achieving a breakthrough level of bonding strength of 210 MPa. This treatment enabled the interface to exhibit a "self-anchoring" effect at high temperatures, with the strength increasing as the temperature rises.

[0124] It should be noted that the above embodiments are only used to illustrate the technical solutions of the present invention rather than to limit them. Although the present invention has been described in detail with reference to the preferred embodiments, those of ordinary skill in the art should understand that the technical solutions of the present invention can be modified or equivalently replaced without departing from the spirit and scope of the technical solutions of the present invention, and all of them should be covered within the scope of the present invention.

Claims

1. A nickel-based superalloy, characterized in that: It includes a base alloy and a cladding alloy; The cladding alloy is a molten solder; The solder is coated on the surface of the base alloy in a molten state to form the nickel-based superalloy; The base alloy has the following components by mass percentage: C ≤ 0.06%; Cr: 16.6 - 17.7%; Nb: 6.1 - 6.8%; N: 6.5 - 7.0%; Ta: 4.3 - 4.8%; Co ≤ 0.7%; Mn: 2.7 - 3.1%; Al: 1.0 - 1.3%; the balance is Fe; The cladding alloy has the following components by mass percentage: Ni: 35 - 45%; Cr: 10 - 15%; Fe: 15 - 20%; W: 3 - 5%; Ti: 1.0 - 1.5%; Rh: 15 - 20%; B: 0.5 - 2.0%; the balance is Zr; Element Nb in the base alloy and element B in the cladding alloy obtain niobium diboride hard phase at the interface; The hard phase is distributed in a network pattern in the interface region and has a thickness of 5 - 20 μm.

2. A preparation method of a nickel-based superalloy, characterized in that: It includes the following steps, Melting the base alloy in an inert atmosphere at a melting temperature of 1250 - 1450 °C for a melting duration of 0.5 - 1.5 h to obtain an alloy ingot blank; Keeping the melted alloy ingot blank at a temperature of 1140 - 1150 °C for 0.5 - 1.5 h, and pressing it into a plate with a press. The thickness of the pressed plate is 1.2 - 1.5 mm; Cooling the compressed plate at a cooling rate of 3 - 5 °C / min to 650 °C, and then cooling it at a cooling rate of 1 - 2 °C / min to 200 - 300 °C to obtain an alloy plate; After coating the cladding alloy on the surface of the alloy plate, the nickel-based superalloy is obtained.

3. The preparation method of the nickel-based superalloy according to claim 2, wherein: The preparation method of the base alloy includes, Weighing each raw material with C ≤ 0.06%, Cr: 16.6 - 17.7%, Nb: 6.1 - 6.8%, N: 6.5 - 7.0%, Ta: 4.3 - 4.8%, Co ≤ 0.7%, Mn: 2.7 - 3.1%, Al: 1.0 - 1.3%, and the balance being Fe by mass percentage; Dividing each raw material into two groups. The first group of raw materials includes Fe, C, and Cr, and the second group includes Nb, N, Ta, Co, Mn, and Al; In a vacuum induction melting furnace, heating the first group of raw materials to 1400 - 1450 °C and holding for 20 - 30 min to obtain molten material a; Meanwhile, in another vacuum induction melting furnace, heating the second group of raw materials to 1400 - 1450 °C and holding for 20 - 30 min to obtain molten material b; Under the protection of an inert atmosphere, pouring molten material a into molten material b, stirring evenly and holding for 10 - 20 min to obtain an alloy melt; Pouring the uniformly mixed alloy melt into a water-cooled steel mold lined with a ceramic cylinder at a speed of 200 - 300 mm / min, and standing for 30 - 60 min to obtain the base alloy; Among them, the material of the lined ceramic cylinder is alumina ceramic, and the inner wall thickness is 5 - 10 mm; The water-cooled steel mold is cooled by circulating water.

4. The preparation method of the nickel-based superalloy according to claim 3, characterized in that: The preparation method of the cladding alloy includes, Raw materials of Ni: 35 - 45 wt%, Cr: 10 - 15 wt%, Fe: 15 - 20 wt%, W: 3 - 5 wt%, Ti: 1.0 - 1.5 wt%, Rh: 15 - 20 wt%, B: 0.5 - 2.0 wt%, and the balance Zr, measured by mass percentage; In an electron beam melting furnace, heat the raw materials to 1400 - 1500 °C under pure argon protection and keep them warm for 0.5 - 1.5 h to obtain a molten cladding alloy; Cool the molten cladding alloy to 800 - 900 °C and maintain the molten state; Using a vacuum magnetron sputtering coating equipment, under the condition that the sputtering power density is 10 - 20 W / cm2, uniformly coat the molten cladding alloy onto the surface of the base alloy plate at a sputtering speed of 150 - 200 s / min; After coating, keep it warm in a vacuum environment at 1000 - 1100 °C for 20 - 30 s to complete the preliminary bonding of the cladding alloy and the base alloy.

5. The preparation method of the nickel-based superalloy according to claim 4, characterized in that: After the preliminary bonding of the cladding alloy and the base alloy, It also includes the following steps, Heat the preliminarily bonded cladding alloy and the base alloy at 2100 - 2400 °C for 2 - 3 h, and then air-cool to 25 °C; Then heat up to 1050 °C and conduct vacuum annealing for 40 - 80 min, and then air-cool to 25 °C to obtain the nickel-based superalloy.