A grain boundary L12 phase-stabilized Ni-Co-Cr-Fe-Al-Ti multi-principal element alloy and its preparation method

Abstract

The present application relates to a kind of grain boundary L12 phase stable Ni-Co-Cr-Fe-Al-Ti multi-main element alloy and its preparation method, belong to the technical field of multi-main element alloy material for high temperature structure.The alloy is based on FCC matrix and the Ni, Co, Cr, Fe, Al, Ti system that can form L12 ordered phase, through the synergistic route of alloy composition design, long time high temperature pre-precipitation treatment before rolling, rolling deformation and subsequent annealing control, 30%~50% volume fraction of L12 phase is formed in advance in alloy, and in subsequent thermal mechanical treatment, the L12 phase is mainly distributed in the form of equiaxed or near equiaxed particles in recrystallized grain boundary or near grain boundary region.The present application realizes the technical effect of grain boundary stable strengthening, the improvement of microstructure thermal stability and the synergistic optimization of strength and plasticity.

Description

Technical Field

[0001] This invention relates to the field of high-temperature structural multi-principal alloy materials, and particularly to a Ni-Co-Cr-Fe-Al-Ti multi-principal alloy that regulates the spatial distribution of the L12 phase through a long-term, medium-high temperature pre-precipitation treatment before rolling. It also relates to a method for preparing a high-temperature structural multi-principal alloy that stably distributes the L12 phase mainly in the form of equiaxed particles at the grain boundaries and suppresses the precipitation of the L12 phase inside the grains. Background Technology

[0002] Multi-principal alloys have attracted widespread attention in the field of high-strength, high-toughness, and high-temperature structural materials due to their large compositional design space, tunable phase structure, and good strengthening and toughening potential. Multi-principal alloys containing Ni, Co, Cr, Fe, Al, and Ti can typically form an FCC matrix and an ordered L12 phase. The L12 phase can improve the alloy strength through its coherent or semi-coherent relationship with the matrix.

[0003] In existing technologies, Ni-Co-Cr-Fe-Al-Ti alloys are mostly strengthened by solution treatment, rolling deformation, and subsequent annealing or aging. Under these process conditions, the L12 phase often forms continuously or discontinuously within the recrystallized grains. For prolonged high-temperature annealing or service environments, high-density lamellar L12 precipitates may also form within the grains. While these precipitates can improve material strength, they are primarily located within the grains, offering limited sustained constraint on grain boundary migration and potentially inducing localized stress concentrations during tensile deformation.

[0004] On the other hand, if a stable second-phase pinning structure is lacking at the grain boundaries, recrystallized grains are prone to local coarsening or even abnormal grain growth under prolonged high-temperature conditions, leading to a decrease in material strength, reduced microstructure uniformity, and decreased service reliability. Therefore, the key to improving the long-term high-temperature microstructure stability of this type of multi-principal element alloy is to transform the L12 phase from an internal precipitate to a grain boundary stable phase mainly distributed at the grain boundaries while maintaining its strengthening potential.

[0005] Existing conventional solution-rolling-annealing routes typically do not introduce high-volume-fraction pre-precipitation of L12 particles before rolling, making it difficult to simultaneously achieve stable pinning of L12 particles at grain boundaries and avoidance of L12 phase precipitation within grains. Therefore, a technical solution that combines compositional design with thermomechanical processing is needed to obtain a stable fine-grained microstructure suitable for long-term, medium-to-high-temperature service environments by controlling the formation sequence, morphology, and spatial distribution of the L12 phase. Summary of the Invention

[0006] In view of this, this invention addresses the problems of grain coarsening and secondary precipitation of L12 phase within grains, leading to poor strength-ductility matching at room temperature and high temperature, which are prone to occur in existing Ni-Co-Cr-Fe-Al-Ti multi-principal alloys during long-term high-temperature annealing or service. It provides a synergistic preparation method based on alloy composition design, long-term high-temperature pre-precipitation treatment before rolling, and control of rolling deformation and annealing. This method ensures that the L12 phase is preferentially and stably distributed at grain boundaries in the form of equiaxed or near-equiaxed particles, and minimizes the precipitation of continuous or discontinuous L12 phase within grains, thereby improving the long-term microstructure stability of the alloy and the strength-ductility matching at room temperature and high temperatures. This preparation method broadens the application range of existing Ni-Co-Cr-Fe-Al-Ti multi-principal alloys under long-term high-temperature service conditions and is suitable for industrial applications.

[0007] To achieve the above objectives, the present invention provides the following solution: One of the technical solutions of this invention is a method for preparing a grain boundary L12 phase-stabilized Ni-Co-Cr-Fe-Al-Ti multi-principal element alloy, characterized by comprising the following steps: Ni, Co, Cr, Fe, Al and Ti raw materials are smelted under an inert gas atmosphere to obtain alloy ingots; The alloy ingot is remelted to form an alloy casting; The alloy casting is homogenized, then cooled in the furnace to the pre-precipitation treatment temperature for pre-precipitation treatment, and then water-quenched to room temperature. The homogenization treatment conditions are set as follows: the temperature is increased from room temperature to 1100-1300℃ at a heating rate of 5-10℃ / min and held for 8-24 hours; The pre-precipitation treatment temperature is 600-900℃, and the pre-precipitation treatment time is 6-24 hours.

[0008] Pre-precipitation treatment can form an L12 phase with a volume fraction of 30% to 50% in the alloy. The L12 phase is equiaxed, cubic or near-equiaxed granular.

[0009] In a preferred embodiment of the present invention, the purity of the Ni, Co, Cr, Fe, Al and Ti raw materials is ≥99.99%; the Ni-Co-Cr-Fe-Al-Ti multi-principal element alloy comprises the following elements in atomic percentage: Ni 35 at.%, Co 30 at.%, Cr 15 at.%, Fe 10 at.%, Al 5 at.%, Ti 5 at.%.

[0010] In a preferred embodiment of the present invention, the melting is electric arc melting, and the melting is carried out 5-10 times and then cooled. The melting current for each melting is 350A and the melting time for each melting is 70s.

[0011] In a preferred embodiment of the present invention, after the pre-precipitation treatment is completed, the process further includes sequential water quenching to room temperature, rolling, and annealing.

[0012] In a preferred embodiment of the present invention, the rolling is carried out by multi-pass cold rolling, with a reduction of 0.15-0.20 mm per pass, and a total thickness reduction of 88-91%.

[0013] In a preferred embodiment of the present invention, the annealing temperature is 600-900°C and the annealing time is 24-144 hours.

[0014] Annealing can cause the matrix to recrystallize and make the L12 phase mainly distributed in the recrystallization grain boundaries, subgrain boundaries or near grain boundaries, forming a Ni-Co-Cr-Fe-Al-Ti multi-principal element alloy with grain boundary L12 grain pinning grain boundary structure. After annealing, the alloy contains ultrafine recrystallized grains and a small amount of residual incompletely recrystallized grains, wherein the average size of the recrystallized grains is 0.5-2.5 μm.

[0015] After annealing at 800℃ for 72 hours, the recrystallization fraction of the alloy is about 75%, the alloy grains still maintain a fine structure, and there is no obvious local grain coarsening or abnormal grain growth. The average recrystallized grain size is about 1.38 μm.

[0016] After annealing, the L12 phase in the alloy is distributed in the form of equiaxed particles at the grain boundaries, and no obvious L12 phase precipitation was observed inside the recrystallized grains.

[0017] The second technical solution of the present invention is a Ni-Co-Cr-Fe-Al-Ti multi-principal element alloy prepared by the above-mentioned preparation method.

[0018] In a preferred embodiment of the present invention, the L12 phase is mainly distributed in the recrystallization grain boundaries, subgrain boundaries or near grain boundaries, forming a grain boundary structure with L12 grains pinning the grain boundaries.

[0019] The third technical solution of this invention is the application of the above-mentioned Ni-Co-Cr-Fe-Al-Ti multi-principal-element alloy in the preparation of aero-engine components, gas turbine components, or energy and power equipment components.

[0020] Traditional Ni-Co-Cr-Fe-Al-Ti multi-principal element alloys typically employ homogenization treatment, rolling deformation, and subsequent annealing or aging processes to achieve a strengthened microstructure. Under these process conditions, the L12 phase often forms within the recrystallized grains in a continuous or discontinuous precipitation manner. Although intragranular L12 precipitates can improve material strength, their primary distribution within the grains limits their sustained constraint on grain boundary migration. Furthermore, during prolonged high-temperature service, the formation of numerous lamellar L12 phases within the grains can lead to inhomogeneous microstructure evolution and may induce localized stress concentrations, thereby reducing the material's microstructural stability and service reliability.

[0021] This invention does not rely solely on material composition design or a single heat treatment process. Instead, it employs a synergistic process route of "alloy composition design, long-term medium-high temperature pre-precipitation, rolling deformation, and annealing control" to regulate the formation sequence, precipitation path, and spatial distribution of the L12 phase. By introducing high-volume-fraction pre-precipitated L12 particles before rolling, L12-forming elements such as Ni, Al, and Ti preferentially participate in the formation of the ordered phase, thereby reducing the chemical driving force for secondary precipitation of the L12 phase within the grains during subsequent annealing. Simultaneously, the pre-formed L12 particles gradually reconstruct from traditional intragranular dispersed precipitates during subsequent rolling and annealing as recrystallization progresses, transforming into equiaxed particles mainly distributed near grain boundaries, thus playing the role of grain boundary stabilizing and grain boundary pinning phases.

[0022] Specifically, this invention introduces approximately 44 vol.% L12 phase before rolling through a pre-precipitation process, allowing it to exist stably in the microstructure as near-equiaxed particles. These particles preferentially accumulate at grain boundaries, subgrain boundaries, and adjacent grain boundary regions during recrystallization, forming continuous or discontinuous grain boundary pinning structures. This effectively restricts grain boundary migration and inhibits local grain coarsening and abnormal grain growth. Simultaneously, since the pre-precipitation stage consumes a large amount of L12-forming elements such as Ni, Al, and Ti, it effectively avoids the excessive formation of lamellar L12 phase within the recrystallized grains, maintaining a microstructure dominated by an FCC matrix within the recrystallized grains.

[0023] Experimental results show that after final annealing, i.e., isothermal aging at 800℃ for 72 hours, the alloy still maintains a bimodal microstructure composed of ultrafine recrystallized grains and residual unrecrystallized grains, with a recrystallization fraction of approximately 75% and an average equivalent circle diameter of approximately 1.38 μm for the recrystallized grains, which is significantly lower than the average recrystallized grain size of similar alloys after the same aging treatment. Furthermore, the Ni-Co-Cr-Fe-Al-Ti multi-principal-element alloy prepared by this invention exhibits excellent strength-ductility matching at room temperature and high temperature. After long-term annealing at 800℃, the pre-precipitated sample achieves a yield strength of approximately 1392 MPa and a total elongation of approximately 19.7% at room temperature; and a yield strength of approximately 931 MPa and a total elongation of approximately 28.4% at 600℃. These combined properties demonstrate that this invention not only maintains excellent strength levels but also possesses good ductility and microstructural stability.

[0024] Compared with the prior art, the present invention has the following beneficial effects: The preparation method of this invention can reduce the supersaturation of L12-forming elements in the matrix, suppress the secondary precipitation of continuous or discontinuous L12 phases inside the grains after annealing, and construct a grain boundary L12 particle pinning structure. This maintains a fine grain structure under long-term, medium-high temperature annealing or service conditions, and inhibits local grain coarsening and abnormal grain growth. This invention effectively solves the problems of easy grain coarsening, difficulty in L12 phase control, and insufficient microstructure stability in existing Ni-Co-Cr-Fe-Al-Ti multi-principal alloys during long-term, medium-high temperature service. It achieves the technical effects of grain boundary stabilization and strengthening, improved microstructure thermal stability, and synergistic optimization of strength and plasticity. This expands the application prospects of this type of multi-principal alloy in aero-engines, gas turbines, energy and power equipment, and other structural components that require long-term, medium-high temperature service, and has good industrial application value. Attached Figure Description

[0025] To more clearly illustrate the technical solutions in the embodiments of the present invention or the prior art, the drawings used in the embodiments will be briefly introduced below. Obviously, the drawings described below are only some embodiments of the present invention. For those skilled in the art, other drawings can be obtained based on these drawings without creative effort.

[0026] Figure 1 XRD patterns of samples prepared in Examples 1-2 and Comparative Examples 1-2; Figure 2 The images shown are scanning electron microscope (SEM) images of Example 1; where (a) is the microstructure of the alloy casting prepared in Example 1; and (b)-(c) are magnified images of the alloy casting prepared in Example 1 and local EDS surface scan results. Figure 3Scanning electron microscope (SEM) images of the sample prepared in Comparative Example 1; where (a) is the microstructure of the alloy casting prepared in Comparative Example 1; and (b) is a magnified partial image of the alloy casting prepared in Comparative Example 1. Figure 4 The images are EBSD diagrams of the samples prepared in Example 2 and Comparative Example 2; where (a) is the inverse pole figure of the alloy plate prepared in Example 2; (b) is the superimposed diagram of the grain boundaries and phase structure of the alloy plate prepared in Example 2; (c) is the inverse pole figure of the alloy plate prepared in Comparative Example 2; and (d) is the superimposed diagram of the grain boundaries and phase structure of the alloy plate prepared in Comparative Example 2. Figure 5 The images show the recrystallized grain size distribution of the samples prepared in Example 2 and Comparative Example 2, obtained by EBSD analysis; where (a) is Example 2 and (b) is Comparative Example 2. Figure 6 The images show magnified scanning electron microscope (SEM) images and local EDS surface scan results of the recrystallized grains of the samples prepared in Example 2 and Comparative Example 2, respectively; where (a)-(b) are from Example 2 and (c)-(d) are from Comparative Example 2. Figure 7 The tensile stress-strain curves of Example 2 and Comparative Example 2 under room temperature and 600°C conditions are shown; where (a) is the room temperature condition and (b) is the 600°C condition. Detailed Implementation

[0027] Various exemplary embodiments of the present invention will now be described in detail. This detailed description should not be considered as a limitation of the present invention, but rather as a more detailed description of certain aspects, features, and embodiments of the present invention.

[0028] It should be understood that the terminology used in this invention is merely for describing particular embodiments and is not intended to limit the invention. Furthermore, with respect to numerical ranges in this invention, it should be understood that each intermediate value between the upper and lower limits of the range is also specifically disclosed. Any stated value or intermediate value within a stated range, as well as each smaller range between any other stated value or intermediate value within said range, is also included in this invention. The upper and lower limits of these smaller ranges may be independently included or excluded from the range.

[0029] Unless otherwise stated, all technical and scientific terms used herein have the same meaning as commonly understood by one of ordinary skill in the art. While only preferred methods and materials have been described herein, any methods and materials similar or equivalent to those described herein may be used in the implementation or testing of this invention. All references to this specification are incorporated by way of citation to disclose and describe methods and / or materials associated with those references. In the event of any conflict with any incorporated reference, the content of this specification shall prevail.

[0030] Various modifications and variations can be made to the specific embodiments described in this specification without departing from the scope or spirit of the invention, as will be apparent to those skilled in the art. Other embodiments derived from this specification will also be obvious to those skilled in the art. This specification and embodiments are merely exemplary.

[0031] The terms “include,” “including,” “have,” “contain,” etc., used in this article are all open-ended terms, meaning that they include but are not limited to.

[0032] This invention provides a grain boundary L12 phase-stabilized Ni-Co-Cr-Fe-Al-Ti multi-principal-element alloy and its preparation method.

[0033] Preferably, the Ni-Co-Cr-Fe-Al-Ti multi-principal-element alloy is composed of the following atomic percentages: Ni 35 at.%, Co 30 at.%, Cr 15 at.%, Fe 10 at.%, Al 5 at.%, and Ti 5 at.%. Ni, Co, and Fe are used to form the FCC matrix and regulate its stability; Cr is used to improve solid solution strengthening and high-temperature microstructure stability; Al and Ti, as L12 ordered phase forming elements, work with Ni to promote L12 phase formation.

[0034] The preparation method of the Ni-Co-Cr-Fe-Al-Ti multi-principal-element alloy includes the following steps: (1) Weigh out the bulk Ni, Co, Cr, Fe, Al and Ti with a purity of not less than 99.99 wt.% according to atomic percentage as raw materials.

[0035] Preferably, the mass of each elemental metal is calculated and weighed based on a total mass of 30 g. Before weighing, the surface of each elemental metal raw material is mechanically ground to remove the surface oxide layer. Subsequently, each raw material is ultrasonically cleaned in anhydrous ethanol and then air-dried for later use.

[0036] Preferably, the weighing is performed using an electronic balance with an accuracy of 0.0001 g, and the weighing error of each raw material is controlled within ±0.0020 g.

[0037] (2) The metal elements processed in step (1) are placed sequentially in a water-cooled copper crucible of an electric arc melting furnace, with the high-melting-point metals Cr and Ti placed on the upper layer and the low-melting-point metal Al placed on the lower layer. The electric arc melting is carried out under the protection of high-purity argon gas, and the alloy button ingots with uniform composition are obtained by repeated turning and remelting.

[0038] Preferably, the electric arc furnace is first evacuated to a vacuum level of 2.0 × 10⁻⁶. -3The pressure is below 1 Pa, and then high-purity argon gas with a purity of not less than 99.999% is introduced as a protective atmosphere.

[0039] Preferably, each alloy button ingot is turned over and melted no less than 7 times.

[0040] Preferably, the mass of a single alloy button ingot is 30 g.

[0041] (3) The alloy button ingot obtained in step (2) is remelted and formed by suction casting using a vacuum arc melting furnace with a suction casting system. The resulting casting is used as the initial blank for subsequent heat treatment and plastic deformation processing.

[0042] Preferably, the molten alloy liquid is drawn into a water-cooled copper mold under negative pressure and solidifies rapidly to obtain a block alloy casting with dimensions of 10 mm × 10 mm × 45 mm.

[0043] (4) Take the Ni obtained in step (3) 35 Co 30 Cr 15 Fe 10 The Al5Ti5 alloy castings undergo homogenization and pre-precipitation treatment.

[0044] Preferably, the homogenization and pre-precipitation treatment in step (4) requires sealing the sample in a quartz tube filled with high-purity inert gas at 0.5 atmospheres, heating it from room temperature to 1100-1300℃ (preferably 1200℃) at a heating rate of about 5-10℃ / min, holding it at that temperature for 8-24 hours (preferably 24 hours), then cooling it in the furnace to 600-900℃ (preferably 800℃), and continuing to hold it at that temperature for 6-24 hours (preferably 10 hours), and then rapidly cooling it to room temperature by water quenching to obtain the pre-precipitated sample.

[0045] The reason why the pre-precipitation treatment temperature is limited to 600-900℃ in this invention is that the FCC+L12 precipitate phase structure is obtained at 600-900℃, while avoiding compositional segregation and the precipitation of irrelevant impurity phases. Only when the parameters are within 600-900℃ does the prepared alloy exhibit a better strength-plasticity match.

[0046] Preferably, the cooling rate of the furnace cooling process is 10-15℃ / min.

[0047] Furthermore, after pre-precipitation treatment, an L12 phase with a volume fraction of 30-50 vol.% is formed inside the alloy.

[0048] Preferably, after pre-precipitation treatment, the volume fraction of the L12 phase in the alloy is 44 vol.%.

[0049] Furthermore, the L12 phase exhibits an equiaxed, nearly equiaxed, or cubic granular morphology.

[0050] Preferably, the average equivalent circular diameter of the L12 particles is 150-400 nm.

[0051] More preferably, the average equivalent circle diameter of the L12 particles is 209 nm.

[0052] (5) The pre-precipitated sample obtained in step 4 is subjected to cold rolling deformation treatment. The rolling direction is along the length of the sample. After multiple cold rolling passes, an alloy plate with a thickness of about 1 mm is obtained.

[0053] Preferably, the reduction per pass is 0.15-0.20 mm.

[0054] Preferably, the total thickness reduction reaches 88-91%.

[0055] (6) The alloy plate obtained in step 5 is sealed in a quartz tube filled with 0.5 atmospheres of high-purity argon gas and placed in a heat treatment furnace preheated to 800°C for annealing. The holding time is 72 hours. After annealing, it is immediately cooled to room temperature by water quenching.

[0056] Unless otherwise specified, the technical solutions described in this invention are all conventional solutions in the field, and the reagents or raw materials used are all purchased from commercial channels or are publicly available unless otherwise specified.

[0057] The technical solutions provided by the present invention will be described in detail below with reference to the embodiments, but they should not be construed as limiting the scope of protection of the present invention.

[0058] Example 1 (T1) A multi-principal alloy ingot containing Ni, Co, Cr, Fe, Al, and Ti. The alloy composition of the multi-principal alloy ingot is Ni 35 at.%, Co 30 at.%, Cr 15 at.%, Fe 10 at.%, Al 5 at.%, and Ti 5 at.%. This multi-principal alloy is smelted from blocky Ni, Co, Cr, Fe, Al, and Ti with a purity of ≥99.99%.

[0059] The method for synthesizing the above-mentioned multi-principal element alloy castings includes the following steps: S1. Batching: The surfaces of each elemental metal are treated to remove the oxide layer, and each raw material is ultrasonically cleaned with anhydrous ethanol cleaning solution. Then, based on the atomic ratios of each elemental metal (Ni 35 at.%, Co 30 at.%, Cr 15 at.%, Fe 10 at.%, Al 5 at.%, Ti 5 at.%), the mass of each elemental metal is calculated and weighed with a total mass of 30 g. Preparation of S2 alloy ingots: Using the electric arc melting method, the various metal elements prepared in S1 were sequentially placed in a water-cooled copper crucible within an electric arc furnace. The electric arc furnace was then evacuated to a vacuum of 2.0 × 10⁻⁶. -3 Pa, then inert gas (argon with a purity ≥ 99.999%) is introduced to 0.5 atmospheres, and the material is repeatedly turned and melted 7 times under the protection of high-purity inert gas to obtain an alloy ingot with uniform composition; the repeated turning and melting of the raw material is carried out using a high-frequency electric arc; the melting current for each melting is 350A, and the melting time for each melting is 70s. S3. Preparation of alloy castings: The alloy button ingots obtained in S2 are remelted in an electric arc furnace, and the molten alloy liquid is drawn into the matching water-cooled copper mold under negative pressure and solidified rapidly to obtain a block alloy casting with dimensions of 10 mm × 10 mm × 45 mm.

[0060] S4. Heat treatment of alloy castings: The Ni obtained in S3... 35 Co 30 Cr 15 Fe 10 The Al5Ti5 alloy casting was sealed in a quartz tube filled with high-purity inert gas (argon with a purity of ≥99.999%) at 0.5 atmospheres. The temperature was increased from room temperature to 1200°C at a heating rate of about 10°C / min. After holding at this temperature for 24 hours, the casting was cooled to 800°C in the furnace and held at this temperature for another 10 hours. Finally, it was rapidly cooled to room temperature by water quenching.

[0061] Example 2 (T2) A multi-principal alloy ingot containing Ni, Co, Cr, Fe, Al, and Ti. The alloy composition of the multi-principal alloy ingot is Ni 35 at.%, Co 30 at.%, Cr 15 at.%, Fe 10 at.%, Al 5 at.%, and Ti 5 at.%. This multi-principal alloy is smelted from blocky Ni, Co, Cr, Fe, Al, and Ti with a purity of ≥99.99%.

[0062] The method for synthesizing the above-mentioned multi-principal element alloy plates includes the following steps: S1. Batching: The surfaces of each elemental metal are treated to remove the oxide layer, and each raw material is ultrasonically cleaned with anhydrous ethanol cleaning solution. Then, based on the atomic ratios of each elemental metal (Ni 35 at.%, Co 30 at.%, Cr 15 at.%, Fe 10 at.%, Al 5 at.%, Ti 5 at.%), the mass of each elemental metal is calculated and weighed with a total mass of 30 g. Preparation of S2 alloy ingots: Using the electric arc melting method, the various metal elements prepared in S1 were sequentially placed in a water-cooled copper crucible within an electric arc furnace. The electric arc furnace was then evacuated to a vacuum of 2.0 × 10⁻⁶. -3 Pa, then inert gas is introduced to 0.5 atmospheres, and the material is repeatedly turned and melted 7 times under the protection of high-purity inert gas to obtain an alloy ingot with uniform composition; the repeated turning and melting of the raw material is carried out using a high-frequency electric arc; the melting current for each melting is 350A, and the melting time for each melting is 70s. S3. Preparation of alloy sheet: The alloy button ingots obtained in S2 are remelted in an electric arc furnace, and the molten alloy liquid is drawn into the matching water-cooled copper mold under negative pressure and solidified rapidly to obtain a block alloy casting with dimensions of 10 mm × 10 mm × 45 mm.

[0063] S4. Heat treatment of alloy castings: The Ni obtained in S3... 35 Co 30 Cr 15 Fe 10 The Al5Ti5 alloy casting was sealed in a quartz tube filled with high-purity inert gas at 0.5 atmospheres. The temperature was increased from room temperature to 1200°C at a rate of about 10°C / min. After holding at this temperature for 24 hours, the casting was cooled to 800°C in the furnace and held at this temperature for another 10 hours. Finally, it was rapidly cooled to room temperature by water quenching.

[0064] S5. Rolling and Annealing of Alloy Castings: The alloy castings obtained in S4 were subjected to multi-pass rolling at room temperature, with the rolling direction along the length of the alloy sample. The reduction per pass was 0.15-0.20 mm, and the total reduction reached 88-91%. Subsequently, the samples were sealed in a quartz tube filled with high-purity argon gas at 0.5 atmospheres. The quartz tube containing the samples was placed in a heat treatment furnace preheated to 800℃ for 72 hours of isothermal aging treatment. Finally, the samples were rapidly cooled to room temperature by water quenching. (That is, the only difference from Example 1 is that after step S4, step S5, rolling and annealing of the alloy casting is added: the alloy casting obtained in S4 is rolled in multiple passes at room temperature, with the rolling direction along the length of the alloy sample, and the reduction in each pass is 0.15-0.20 mm, and the total reduction in rolling needs to reach 88-91%. Then the sample is sealed in a quartz tube filled with high-purity argon at 0.5 atmospheres, and the quartz tube containing the sample is placed in a heat treatment furnace preheated to 800°C for 72 hours of isothermal aging treatment, and finally rapidly cooled to room temperature by water quenching.) Comparative Example 1 (T3) A multi-principal alloy ingot containing Ni, Co, Cr, Fe, Al, and Ti. The alloy composition of the multi-principal alloy ingot is Ni 35 at.%, Co 30 at.%, Cr 15 at.%, Fe 10 at.%, Al 5 at.%, and Ti 5 at.%. This multi-principal alloy is smelted from blocky Ni, Co, Cr, Fe, Al, and Ti with a purity of ≥99.99%.

[0065] The method for synthesizing the above-mentioned multi-principal element alloy castings includes the following steps: S1. Batching: The surfaces of each elemental metal are treated to remove the oxide layer, and each raw material is ultrasonically cleaned with anhydrous ethanol cleaning solution. Then, based on the atomic ratios of each elemental metal (Ni 35 at.%, Co 30 at.%, Cr 15 at.%, Fe 10 at.%, Al 5 at.%, Ti 5 at.%), the mass of each elemental metal is calculated and weighed with a total mass of 30 g. Preparation of S2 alloy ingots: Using the electric arc melting method, the various metal elements prepared in S1 were sequentially placed in a water-cooled copper crucible within an electric arc furnace. The electric arc furnace was then evacuated to a vacuum of 2.0 × 10⁻⁶. -3 Pa, then inert gas is introduced to 0.5 atmospheres, and the material is repeatedly turned and melted 7 times under the protection of high-purity inert gas to obtain an alloy ingot with uniform composition; the repeated turning and melting of the raw material is carried out using a high-frequency electric arc; the melting current for each melting is 350A, and the melting time for each melting is 70s. S3. Preparation of alloy castings: The alloy button ingots obtained in S2 are remelted in an electric arc furnace, and the molten alloy liquid is drawn into the matching water-cooled copper mold under negative pressure and solidified rapidly to obtain a block alloy casting with dimensions of 10 mm × 10 mm × 45 mm.

[0066] S4. Heat treatment of alloy castings: The Ni obtained in S3... 35 Co 30 Cr 15 Fe 10 The Al5Ti5 alloy casting was sealed in a quartz tube filled with high-purity inert gas at 0.5 atmospheres. The temperature was increased from room temperature to 1200°C at a rate of approximately 10°C / min, held at that temperature for 24 hours, and then rapidly cooled to room temperature by water quenching. (The only difference from Example 1 is that the step of "cooling to 800°C in the furnace and holding for another 10 hours" is omitted in step S4.) Comparative Example 2 (T4) A multi-principal alloy ingot containing Ni, Co, Cr, Fe, Al, and Ti. The alloy composition of the multi-principal alloy ingot is Ni 35 at.%, Co 30 at.%, Cr 15 at.%, Fe 10 at.%, Al 5 at.%, and Ti 5 at.%. This multi-principal alloy is smelted from blocky Ni, Co, Cr, Fe, Al, and Ti with a purity of ≥99.99%.

[0067] The method for synthesizing the above-mentioned multi-principal element alloy plates includes the following steps: S1. Batching: The surfaces of each elemental metal are treated to remove the oxide layer, and each raw material is ultrasonically cleaned with anhydrous ethanol cleaning solution. Then, based on the atomic ratios of each elemental metal (Ni 35 at.%, Co 30 at.%, Cr 15 at.%, Fe 10 at.%, Al 5 at.%, Ti 5 at.%), the mass of each elemental metal is calculated and weighed with a total mass of 30 g. Preparation of S2 alloy ingots: Using the electric arc melting method, the various metal elements prepared in S1 were sequentially placed in a water-cooled copper crucible within an electric arc furnace. The electric arc furnace was then evacuated to a vacuum of 2.0 × 10⁻⁶. -3 Pa, then inert gas is introduced to 0.5 atmospheres, and the material is repeatedly turned and melted 7 times under the protection of high-purity inert gas to obtain an alloy ingot with uniform composition; the repeated turning and melting of the raw material is carried out using a high-frequency electric arc; the melting current for each melting is 350A, and the melting time for each melting is 70s. S3. Preparation of alloy castings: The alloy button ingots obtained in S2 are remelted in an electric arc furnace, and the molten alloy liquid is drawn into the matching water-cooled copper mold under negative pressure and solidified rapidly to obtain a block alloy casting with dimensions of 10 mm × 10 mm × 45 mm.

[0068] S4. Heat treatment of alloy castings: The Ni obtained in S3... 35 Co 30 Cr 15 Fe 10 The Al5Ti5 alloy casting was sealed in a quartz tube filled with high-purity inert gas at 0.5 atmospheres and heated from room temperature to 1200°C at a heating rate of about 10°C / min. After holding at that temperature for 24 hours, it was rapidly cooled to room temperature by water quenching.

[0069] S5. Rolling and Annealing of the Alloy Casting: The alloy casting obtained in S4 is rolled in multiple passes at room temperature along the length of the alloy sample. Each pass reduces the alloy casting by 0.15-0.20 mm, and the total reduction is 88-91%. The sample is then sealed in a quartz tube filled with 0.5 atmospheres of high-purity argon. The quartz tube containing the sample is placed in a heat treatment furnace preheated to 800°C for 72 hours of isothermal aging. Finally, it is rapidly cooled to room temperature by water quenching. (That is, the only difference from Example 2 is that the step of "cooling to 800°C in the furnace and holding for 10 hours" is omitted in step S4.) The samples obtained in Example 2 and Comparative Example 2 were subjected to tensile tests at room temperature and 600°C. Figure 7 The figures show the tensile stress-strain curves of Example 2 and Comparative Example 2 under room temperature and 600°C high temperature conditions. The sample of Example 2 is labeled T2, and the sample of Comparative Example 2 is labeled T4. Table 1 shows the room temperature and high temperature tensile test results of samples T2 and T4.

[0070] Table 1. Results of room temperature and high temperature tensile tests for samples T2 and T4

[0071] As shown in Table 1, the pre-precipitated T2 sample exhibits superior mechanical properties compared to the untreated T4 sample after the same thermomechanical treatment. The room temperature yield strength of the T4 sample is 1053 MPa, the tensile strength reaches 1376 MPa, and the elongation at break is 12.1%. In contrast, the pre-precipitated T2 sample shows an increased room temperature yield strength of 1392 MPa, a tensile strength of 1506 MPa, and an elongation at break approaching 20%. The yield strength of the alloy sample T4 is increased by over 30%, and the elongation at break is as high as 63%. Based on the XRD results of samples T1-T4 (… Figure 1 Ni prepared using this method 35 Co 30 Cr 15 Fe 10 All Al5Ti5 alloys exhibited an FCC structure containing the L12 phase. SEM and EDS surface scan results before rolling showed that the pre-precipitated T1 sample contained an L12 phase with a volume fraction of approximately 44 vol.% and an average equivalent circle diameter of approximately 209 nm. Figure 2 However, no obvious L12 phase formation was observed in the T3 sample, which only underwent homogenization treatment. Figure 3 Furthermore, the EBSD results for samples T2 and T4 indicate that ( Figure 4The alloy, after pre-precipitation treatment and isothermal aging at 800℃ for 72 hours, yielded recrystallized grains with finer average sizes, while no obvious local coarsening effect was observed. This is in conjunction with the grain size distribution of samples T2 and T4 (…). Figure 5 Comparative analysis revealed that the recrystallized grain size of the T2 sample after pre-precipitation treatment was reduced by approximately 42% compared to the T4 sample without pre-precipitation treatment. This demonstrates that pre-precipitation treatment can effectively suppress local coarsening and growth of recrystallized grains under long-term high-temperature service conditions. More importantly, compared to the T4 sample, which exhibits the traditional L12 phase precipitation pattern and a large number of lamellar L12 phases forming within the recrystallized grains, this invention alters the precipitation path and spatial distribution of the L12 phase through the pre-precipitation process. This results in the T2 sample acquiring a microstructure characterized by an intragranular FCC matrix and L12-type equiaxed particles enriched at grain boundaries. Figure 6 This microstructure effectively avoids the formation of large amounts of lamellar L12 phase within the grains, while utilizing the pinning effect of L12 particles at the grain boundaries to inhibit grain boundary migration and grain growth, thereby significantly improving the microstructure stability of the alloy under long-term high-temperature service conditions. Figure 6 Mechanical tensile tests at room temperature and high temperature conditions show that this special microstructure contributes to the Ni... 35 Co 30 Cr 15 Fe 10 Al5Ti5 alloys exhibit superior comprehensive mechanical properties under both room temperature and high temperature conditions. Figure 7 ).

[0072] The various embodiments described in this specification are presented in a progressive manner, with each embodiment focusing on its differences from other embodiments. Similar or identical parts between embodiments can be referred to interchangeably. The above description of the disclosed embodiments enables those skilled in the art to make or use the invention. Various modifications to these embodiments will be readily apparent to those skilled in the art, and the general principles defined herein may be implemented in other embodiments without departing from the spirit or scope of the invention. Therefore, the invention is not to be limited to the embodiments shown herein, but is to be accorded the widest scope consistent with the principles and novel features disclosed herein.

Claims

1. A method for preparing a grain boundary L12 phase-stabilized Ni-Co-Cr-Fe-Al-Ti multi-principal element alloy, characterized in that, Includes the following steps: Ni, Co, Cr, Fe, Al and Ti raw materials are smelted under an inert gas atmosphere to obtain alloy ingots; The alloy ingot is remelted to form an alloy casting; The alloy casting is homogenized, then cooled in the furnace to the pre-precipitation treatment temperature for pre-precipitation treatment, and then water-quenched to room temperature. The homogenization treatment conditions are set as follows: the temperature is increased from room temperature to 1100-1300℃ at a heating rate of 5-10℃ / min and held for 8-24 hours; The pre-precipitation treatment temperature is 600-900℃, and the pre-precipitation treatment time is 6-24 hours.

2. The preparation method according to claim 1, characterized in that, The purity of the Ni, Co, Cr, Fe, Al and Ti raw materials is ≥99.99%; the Ni-Co-Cr-Fe-Al-Ti multi-principal element alloy includes the following atomic percentages of elements: Ni 35 at.%, Co 30 at.%, Cr 15 at.%, Fe 10 at.%, Al 5 at.%, Ti 5 at.%.

3. The preparation method according to claim 1, characterized in that, The melting process is an electric arc melting process, which is carried out 5-10 times and then cooled. The melting current for each melting process is 350A and the melting time for each melting process is 70s.

4. The preparation method according to claim 1, characterized in that, The pre-precipitation treatment is followed by a series of steps including water quenching to room temperature, rolling, and annealing.

5. The preparation method according to claim 4, characterized in that, The rolling process employs a multi-pass cold rolling method, with each pass reducing the thickness by 0.15-0.20 mm, resulting in a total thickness reduction of 88%-91%.

6. The preparation method according to claim 4, characterized in that, The annealing temperature is 600-900℃, and the annealing time is 24-144 hours.

7. A Ni-Co-Cr-Fe-Al-Ti multi-principal-element alloy prepared by the preparation method according to any one of claims 1 to 6.

8. The Ni-Co-Cr-Fe-Al-Ti multi-principal-element alloy according to claim 7, characterized in that, The L12 phase is mainly distributed in recrystallization grain boundaries, subgrain boundaries, or near grain boundaries, forming a grain boundary structure with L12 grains pinning the grain boundaries.

9. The application of the Ni-Co-Cr-Fe-Al-Ti multi-principal-element alloy as described in claim 7 or 8 in the preparation of aero-engine components, gas turbine components or energy and power equipment components.

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