Fe-Co-Ni-Al series high-entropy alloy with extremely high work hardening capacity and preparation method of Fe-Co-Ni-Al series high-entropy alloy

Abstract

The invention discloses a Fe-Co-Ni-Al series high-entropy alloy with extremely high work hardening capacity and a preparation method of the Fe-Co-Ni-Al series high-entropy alloy. A, b, c, d, e, f, g and h all represent atomic percentages corresponding to the elements, 40% < = a < = 50%, 10% < = b < = 20%, 20% < = c < = 30%, 5% < = d < = 15%, 0.4% < = e < = 2%, 0.4% < = f < = 2%, 0.4% < = g < = 2%, 0.01% < = h < = 1%, and a + b + c + d + e + f + g + h = 100%. The high-entropy alloy has good comprehensive mechanical properties, the yield strength of the high-entropy alloy can exceed 1 GPa, the tensile strength of the high-entropy alloy is not lower than 2 GPa, the plastic deformation amount of the high-entropy alloy can reach 20% or above, the work hardening rate of the high-entropy alloy in the plastic deformation stage can reach 30 GPa to the maximum, the high-entropy alloy shows extremely excellent work hardening capacity, brittle fracture under the high-load working condition can be effectively avoided, and the high-entropy alloy can be used for manufacturing a high-strength steel plate. And the performance requirements in a specific working environment are met.

Description

Technical Field

[0001] The present invention relates to an Fe-Co-Ni-Al high-entropy alloy, and particularly to an Fe-Co-Ni-Al high-entropy alloy with extremely high work hardening ability and a preparation method thereof, belonging to the technical field of metallic materials. Background Art

[0002] In recent years, alloys with ultra-high strength have played an indispensable role in many fields, including aerospace, transportation, etc. Existing methods for improving the strength of alloys include introducing precipitation phases and refining grains, etc. However, while these methods improve strength, they also cause a loss of plasticity, resulting in the alloy being prone to brittle fracture under high-load working environments, greatly reducing the service safety of the alloy as a structural material. To enable the alloy to have good plasticity on the premise of ultra-high strength, it is usually necessary to improve the work hardening ability of the alloy during the plastic deformation stage. Therefore, how to effectively improve the work hardening ability of the alloy so as to achieve the improvement of strength-plasticity and service stability has become an urgent problem to be solved.

[0003] High-entropy alloys formed by mixing five or more metal elements in equiatomic or near-equiatomic ratios have broken through the traditional alloy design concept, opening up a wide composition design space for designing alloys with excellent properties. At present, some studies have focused on using the characteristics of high-entropy alloys to improve the work hardening ability. There are mainly two strategies for improving the work hardening ability of high-entropy alloys - constructing heterogeneous structures and introducing transformation-induced plasticity (TRIP) effects. The heterogeneous structure in the alloy refers to a structure composed of two or more regions or phases with different properties, including bimodal structure, gradient structure, and layered heterogeneous structure, etc. Introducing a heterogeneous structure in the alloy can effectively increase the geometrically necessary dislocation density accumulated at the heterogeneous interface during the plastic deformation process, thereby generating back stress and forward stress. While improving the alloy's ability to store dislocations, the movement of dislocations is hindered by dislocation interactions, and finally excellent work hardening ability is obtained. This strengthening method is called heterogeneous deformation-induced (HDI) strengthening. On the other hand, a large number of studies have shown that by regulating the alloy composition, the stability of austenite in the alloy can be effectively reduced, causing martensitic transformation during the plastic deformation process. Since the occurrence of martensitic transformation absorbs a large amount of deformation energy, the alloy can continue to deform without breaking, which is the TRIP effect. Moreover, the martensite generated by the phase transformation usually has a much higher hardness than austenite, and can hinder the movement of dislocations more effectively than austenite, thereby improving the work hardening ability of the alloy.

[0004] However, the above two measures to improve the work hardening ability still have limitations. For example, the high stress concentration at the hetero - structure interface often promotes the nucleation and growth of micro - cracks, leading to premature fracture of the alloy; and under a large deformation degree, the deformation of the hard phase causes a decrease in the geometrically necessary dislocation density at the hetero - interface, resulting in a decline in the work hardening ability; in addition, a large amount of martensitic transformation will significantly reduce the volume fraction of austenite with good coordinated deformation ability, thus causing the formation of micro - cracks between grains and hindering the further improvement of the strength and plasticity of the alloy. Therefore, how to break through the existing strength - plasticity limit by regulating the hetero - structure characteristics and adjusting the martensitic transformation behavior is still a great challenge. Summary of the Invention

[0005] Aiming at the performance defects of existing structural materials, the first object of the present invention is to provide an Fe - Co - Ni - Al high - entropy alloy with extremely high work hardening ability. The room - temperature yield strength of this high - entropy alloy can reach 1 GPa, the uniform elongation can reach more than 20%, and the work hardening rate in the plastic deformation stage can reach 30 GPa, which has broad application prospects in the fields of aerospace, transportation, etc.

[0006] The second object of the present invention is to provide a preparation method for the Fe - Co - Ni - Al alloy, which has simple process, low cost and is easy to realize industrial production.

[0007] To achieve the above technical objects, the present invention provides an Fe - Co - Ni - Al high - entropy alloy with extremely high work hardening ability, and its general formula of elemental composition is Fe a Co b Ni c Al d Ti e Nb f Ta g B h , where a, b, c, d, e, f, g and h all represent the atomic percentages corresponding to each element, 40% ≤ a ≤ 50%, 10% ≤ b ≤ 20%, 20% ≤ c ≤ 30%, 5% ≤ d ≤ 15%, 0.4% ≤ e ≤ 2%, 0.4% ≤ f ≤ 2%, 0.4% ≤ g ≤ 2%, 0.01% ≤ h ≤ 1%, and a + b + c + d + e + f + g + h = 100%; it has a duplex microstructure of a face - centered cubic crystal structure phase and a B2 ordered body - centered cubic crystal structure phase, and there are nano - scale L12 ordered face - centered cubic crystal structure precipitation phases.

[0008] In the Fe-Co-Ni-Al series high-entropy alloy of the present invention, Fe is a ferromagnetic element with good solid-solution strengthening ability. Ni is a face-centered cubic (FCC) phase-forming element. Co element can help reduce the austenite stability to promote martensitic transformation. The role of Al element is to produce coherent L12 precipitation phases and hard B2 ordered phases with the matrix. Ta, Ti, and Nb are L12 precipitation phase stabilizing elements, and a small amount of B element can refine grains and strengthen grain boundaries.

[0009] The Fe-Co-Ni-Al series alloy provided by the present invention has a metastable face-centered cubic structure austenite, which can effectively provide excellent work hardening ability through martensitic transformation during plastic deformation.

[0010] As a preferred solution, in the general formula of the element composition of the Fe-Co-Ni-Al series alloy, 40% ≤ a ≤ 45%, 13% ≤ b ≤ 20%, 23% ≤ c ≤ 30%, 8% ≤ d ≤ 13%, 0.5% ≤ e ≤ 1%, 0.5% ≤ f ≤ 1%, 0.5% ≤ g ≤ 1%, 0.01% ≤ h ≤ 0.1%, and a + b + c + d + e + f + g + h = 100%. For the most preferred Fe-Co-Ni-Al series alloy, in its general formula of composition, 14% ≤ b ≤ 18%, 24% ≤ c ≤ 28%, 9% ≤ d ≤ 12%. If the Al content is too low, a high-density L12 precipitation phase cannot be formed. If the contents of Ta, Ti, and Nb are too low, the L12 precipitation phase cannot be stabilized. An appropriate ratio of Fe and Ni elements can form an FCC austenite matrix. If the Fe element content is too high and the Ni content is too low, a brittle BCC phase is likely to be formed. However, appropriately increasing the Fe and Co elements or decreasing the Ni element can reduce the austenite stability to promote martensitic transformation. Therefore, the element composition of the high-entropy alloy needs to be controlled within a suitable range.

[0011] The present invention also provides a preparation method of a Fe-Co-Ni-Al series alloy, which is obtained by successively cleaning, vacuum melting, vacuum casting, solution treatment, cold rolling, recrystallization annealing treatment, and aging treatment of Fe, Co, Ni, Al, Ta, Nb, Ti, and B raw materials.

[0012] The preparation method of the Fe-Co-Ni-Al series alloy of the present invention can adopt the traditional melting and casting process of the existing technology, which is easy to realize industrial production.

[0013] The Fe-Co-Ni-Al series alloy of the present invention forms an FCC crystal structure with both blocky B2 ordered phase inlays and L12 nano-ordered phase precipitates after cold rolling, recrystallization annealing treatment, and aging treatment. The obtained Fe-Co-Ni-Al series alloy has good mechanical properties, with a room temperature yield strength reaching 1.2 GPa, a tensile strength reaching 2.3 GPa, a uniform elongation rate reaching 22%, and a work hardening rate reaching up to 30 GPa.

[0014] As a preferred solution, the cleaning process includes a step of grinding to remove the surface oxide layer and a step of ultrasonic washing. In the grinding process, mechanical means such as sandpaper or a grinding wheel are usually used to remove the surface metal oxide scale of the raw metals Fe, Ni, Al, Ta, and Ti. The ultrasonic washing step mainly uses industrial ethanol for ultrasonic oscillation cleaning of the raw metals to remove surface organic pollutants. Through surface cleaning pretreatment, the introduction of impurity elements into the alloy material can be effectively reduced. The non-metallic element B does not require cleaning treatment.

[0015] As a preferred solution, in the process of vacuum melting: the elements are sorted according to their melting points and element types for the placement order. First, B is placed at the bottom, and the remaining elements are placed in sequence as Al, Ti, Ni, Co, Fe, Nb, and Ta. Among them, Nb and Ta, as high melting point elements, are placed at the top.

[0016] As a preferred solution, the conditions for vacuum melting are: the arc current is 300 - 400 A, the melting time for each time is 80 - 120 s. After each melting, the ingot is flipped and cooled, and then the next melting is carried out. The melting is repeated at least eight times to make the alloy fully melted and uniform. The time for the ingot to be flipped and cooled is more than 300 s.

[0017] As a preferred solution, the conditions for solution treatment are: at a temperature of 1200 - 1300 °C, after holding for 2 - 24 hours, water quenching is carried out. If the solution treatment temperature is too high, it will cause the alloy to remelt, and if the solution temperature is too low, it will result in non-uniform composition. If the solution treatment time is too long, the grain size will grow too large.

[0018] As a preferred solution, the conditions for cold rolling are: the ambient temperature is 25 - 30 °C, the reduction per pass is 3 - 5 mm, and the total rolling reduction is 80% - 90%. If the rolling reduction is too small, the mechanical properties of the alloy will deteriorate, and if the rolling reduction is too large, the sample is prone to bending, resulting in excessive internal stress.

[0019] As a preferred embodiment, the conditions for the recrystallization annealing treatment are as follows: the temperature is 1000 - 1200 °C, and the time is 5 - 15 minutes. Appropriate recrystallization treatment can eliminate the deformation internal stress in the alloy, reduce the dislocation density, and obtain a stable equiaxed crystal structure. If the recrystallization temperature is too high or the recrystallization time is too long, the grain size will be too large, resulting in a decrease in strength. If the recrystallization temperature is too low or the time is too short, the alloy's recovery and recrystallization will be incomplete, leading to a decrease in plasticity.

[0020] As a preferred embodiment, the conditions for the aging treatment are as follows: the temperature is 600 - 800 °C, and the time is 1 - 3 hours. Through aging heat treatment under appropriate conditions, new phases will precipitate in the alloy, and the yield strength and tensile strength of the alloy will be significantly improved, thereby making it have better toughness and ductility. Further preferably, the aging treatment temperature is 650 - 750 °C. Further preferably, the aging treatment time is 2 - 3 hours.

[0021] As a preferred embodiment, the purity of the metal raw materials Fe, Co, Ni, Al, Ta, Nb, and Ti is greater than 99.9%.

[0022] The present invention provides a method for preparing an Fe - Co - Ni - Al - based alloy, comprising the following steps:

[0023] (1) Use sandpaper or a grinding wheel to remove the surface oxide scale of the raw metals Fe, Co, Ni, Al, Ta, Nb, and Ti, and use industrial ethanol for ultrasonic oscillation cleaning of the raw metals, and then dry them.

[0024] (2) Convert the metallurgical raw materials Fe, Co, Ni, Al, Ta, Nb, and Ti elements into mass ratios according to the atomic percentages in the Fe - Co - Ni - Al - based alloy expression for batching, and use a balance with an accuracy of 0.001 g for batching.

[0025] (3) Feed the raw materials prepared in (2) into a vacuum smelting furnace. The placement order is as follows: first place B at the bottom, and the remaining elements are placed in sequence as Al, Ti, Ni, Co, Fe, Nb, and Ta. Among them, Nb and Ta, as high - melting - point elements, need to be placed at the top. Then adjust the vacuum degree of the vacuum chamber of the vacuum smelting furnace to 5×10 -3 Pa, then fill the furnace cavity with argon gas to half an atmospheric pressure, and then turn on the vacuum smelting furnace to conduct vacuum melting treatment on the alloy raw materials. The conditions for vacuum melting are as follows: the arc current is 300 - 400 A, the time for each melting is 80 - 120 s. After each melting ends, turn the ingot over for cooling, and then conduct the next melting. Repeat melting at least eight times to make the alloy fully melted and uniform. After melting, use a vacuum suction casting device to suction - cast the alloy into a water - cooled copper mold to obtain an alloy material with a B2 ordered body - centered cubic structure and a face - centered cubic structure.

[0026] (4)Alloy preparation: The alloy obtained by melting and suction casting in (3) is sealed in a vacuum quartz tube and subjected to high-temperature solution treatment. The treatment conditions are 1200 - 1300 °C, holding for 2 - 24 hours followed by water quenching.

[0027] (5)Cold rolling treatment: The alloy prepared in (4) is polished with 240# sandpaper to remove the oxide layer, and cold rolling is carried out at an ambient temperature of 25 - 30 °C. The reduction per pass is 3 - 5 mm, and the total cold rolling reduction is 80% - 90%.

[0028] (6)Recrystallization annealing treatment: The Fe-Co-Ni-Al series alloy needs to be subjected to recrystallization annealing treatment. The treatment conditions are 1000 - 1200 °C, for 5 - 15 minutes followed by water quenching.

[0029] (7)Aging annealing treatment: The Fe-Co-Ni-Al series alloy is directly subjected to aging annealing after cold rolling treatment. The treatment conditions are 600 - 800 °C, for 1 - 3 hours followed by water quenching, thus obtaining the Fe-Co-Ni-Al series alloy of the present invention with extremely high work hardening ability.

[0030] The key to the technical solution of the present invention lies in improving the work hardening ability of the alloy during the deformation stage through recrystallization annealing and cold rolling aging processes, and obtaining an Fe-Co-Ni-Al series alloy with excellent comprehensive properties of strength and elongation. Compared with the prior art, the beneficial technical effects brought by the technical solution of the present invention are:

[0031] (1)The general composition formula of the Fe-Co-Ni-Al series alloy provided by the present invention is: Fe a Co b Ni c Al d Ti e Nb f Ta g B h , where a, b, c, d, e, f, g, and h all represent the atomic percentages corresponding to each element, 40% ≤ a ≤ 50%, 10% ≤ b ≤ 20%, 20% ≤ c ≤ 30%, 5% ≤ d ≤ 15%, 0.4% ≤ e ≤ 2%, 0.4% ≤ f ≤ 2%, 0.4% ≤ g ≤ 2%, 0.01% ≤ h ≤ 1%, and a + b + c + d + e + f + g + h = 100%. It has a duplex microstructure of a face-centered cubic crystal structure phase and a B2 ordered body-centered cubic crystal structure phase, and there are nano-scale L12 ordered face-centered cubic crystal structure precipitation phases. The room temperature yield strength of this alloy reaches 1.2 GPa, the tensile strength reaches 2.3 GPa, the uniform elongation reaches 22%, and the work hardening rate is up to 30 GPa at most.

[0032] (2) The Fe-Co-Ni-Al alloy provided by the present invention has a metastable face-centered cubic structure austenite, which can effectively provide excellent work hardening ability through martensitic transformation during plastic deformation.

[0033] (3) The Fe-Co-Ni-Al alloy provided by the present invention can achieve a good match between strength and plasticity.

[0034] (4) The preparation method of the Fe-Co-Ni-Al alloy provided by the present invention has a simple process and is easy to realize industrial production. The metal single-element raw materials in the prepared alloy material are evenly mixed and have excellent mechanical properties. At the same time, the Fe-Co-Ni-Al alloy described in the present invention is obtained by melting ordinary pure metal single-element raw materials, with low preparation cost and no expensive metals such as Mo, and has good application prospects. Description of the Drawings

[0035] Figure 1 The work hardening rate-true strain curve, room temperature tensile stress-strain curve, synchrotron radiation diffraction (HEXRD) pattern and scanning electron microscope (SEM) image of the Fe-Co-Ni-Al alloy prepared in Example 1.

[0036] Figure 2 The work hardening rate-true strain curve, room temperature tensile stress-strain curve, synchrotron radiation diffraction (HEXRD) pattern and scanning electron microscope (SEM) image of the Fe-Co-Ni-Al alloy prepared in Example 2.

[0037] Figure 3 The work hardening rate-true strain curve, room temperature tensile stress-strain curve, synchrotron radiation diffraction (HEXRD) pattern and scanning electron microscope (SEM) image of the Fe-Co-Ni-Al alloy prepared in Example 3.

[0038] Figure 4 The work hardening rate-true strain curve, room temperature tensile stress-strain curve, synchrotron radiation diffraction (HEXRD) pattern and scanning electron microscope (SEM) image of the Fe-Co-Ni-Al alloy prepared in Example 4.

[0039] Figure 5 The work hardening rate-true strain curve, room temperature tensile stress-strain curve, synchrotron radiation diffraction (HEXRD) pattern and scanning electron microscope (SEM) image of the Fe-Co-Ni-Al alloy prepared in Example 5.

[0040] Figure 6The work-hardening rate-true strain curve, room-temperature tensile stress-strain curve, synchrotron radiation diffraction (HEXRD) pattern, and scanning electron microscope (SEM) image of the Fe-Co-Ni-Al alloy prepared in Example 6.

[0041] Figure 7 For comparison, the tensile properties of the Fe-Co-Ni-Al alloys prepared in Comparative Examples 1-3 without Ti, Nb, Ta, and B elements.

[0042] Figure 8 For comparison, the tensile property curves and work-hardening curves of the Co recrystallized alloy and Ni recrystallized alloy in the Fe-Co-Ni-Al alloys prepared in Comparative Examples 4-5. Among them, (a) and (b) are the tensile property curves and work-hardening curves of the Co recrystallized alloy; (c) and (d) are the tensile property curves and work-hardening curves of the Ni recrystallized alloy. Detailed implementation manners

[0043] To facilitate the understanding of the present invention, the content of the present invention will be described more comprehensively and in detail below in conjunction with the accompanying drawings of the specification and preferred embodiments. However, the protection scope of the present invention is not limited to the following specific embodiments.

[0044] Unless otherwise defined, all professional terms used hereinafter have the same meaning as commonly understood by those skilled in the art. The professional terms used herein are only for the purpose of describing specific embodiments and are not intended to limit the protection scope of the present invention.

[0045] Unless otherwise specifically stated, various raw materials, reagents, instruments, and equipment used in the present invention can be obtained through market purchases or can be prepared by existing methods. The test methods are all carried out according to industry standards.

[0046] Example 1

[0047] The general formula for the composition of the Fe-Co-Ni-Al alloy is: Fe a Co b Ni c Al d Ti e Nb f Ta g B h , where a, b, c, d, e, f, g are 43.96%, 14%, 28%, 11.5%, 0.83%, 0.83%, 0.83%, 0.05% respectively.

[0048] The preparation method of the Fe-Co-Ni-Al alloy includes the following steps:

[0049] (1) Use sandpaper or a grinding wheel to remove the surface scale of the raw metals Fe, Co, Ni, Al, Ta, Nb, and Ti. Clean the raw metals by ultrasonic oscillation with industrial ethanol, and then dry them. Convert the atomic percentages of the raw elements Fe, Co, Ni, Al, Ta, Ti, Nb, and B into mass ratios according to the expression and weigh the materials. Use a balance with a precision of 0.001 g for batching; Feed the prepared raw materials into a vacuum smelting furnace. Each element needs to be sorted according to the melting point and element type for the order of placing the raw materials. Specifically, place B at the bottom first, and the remaining elements are Al, Ti, Ni, Co, Fe, Nb, Ta in sequence. As high melting point elements, Nb and Ta need to be placed at the top for subsequent melting. Adjust the vacuum degree of the vacuum chamber of the vacuum smelting furnace to 5×10 -3 Pa, then fill the furnace cavity with argon gas to half an atmosphere, and then turn on the vacuum smelting furnace to melt the alloy raw materials. The arc current is 390 A, and repeat the melting 8 times until the alloy is fully melted and homogenized. Each melting time is 90 s; After melting, use a vacuum suction casting device to suck-cast the alloy into a water-cooled copper mold to obtain an Fe-Co-Ni-Al-based alloy containing a B2 ordered body-centered cubic structure phase and a face-centered cubic structure phase;

[0050] (2) Solution treatment: Seal the alloy obtained by melting and suction casting in (1) in a vacuum quartz tube and perform high-temperature solution treatment. The treatment conditions are 1200 °C, hold for 2 hours, and then water quench.

[0051] (3) Cold rolling treatment: Polish the alloy obtained by solution treatment in (2) with 240# sandpaper to remove the oxide layer, and use a rolling mill to perform cold rolling treatment at room temperature. The reduction per pass is 4 mm, and the total rolling reduction is 90%.

[0052] (4) Recrystallization annealing treatment: Put the alloy obtained by cold rolling treatment in (3) into a heat treatment furnace for recrystallization annealing treatment at a temperature of 1150 °C for 10 min, and then water quench to obtain the recrystallization annealed alloy.

[0053] (5) Preparation of recrystallized age-hardened alloy: Perform age-hardening annealing treatment on the alloy obtained by recrystallization annealing treatment in (4) at a temperature of 650 °C for 2 h, and then water quench.

[0054] The obtained Fe-Co-Ni-Al-based alloy is named Co650 alloy.

[0055] Example 2

[0056] The only difference compared with Example 1 is that the aging temperature is 700 °C.

[0057] The obtained Fe-Co-Ni-Al-based alloy is named Co700 alloy.

[0058] Example 3

[0059] Compared with Example 1, the only difference is that the aging temperature is 750 °C.

[0060] The obtained Fe-Co-Ni-Al alloy system is named Co750 alloy.

[0061] Example 4

[0062] Compared with Example 1, the only difference is that for Fe a Co b Ni c Al d Ti e Nb f Ta g B h a, b, c, d, e, f, g are 43.96%, 17%, 25%, 11.5%, 0.83%, 0.83%, 0.83%, 0.05% respectively.

[0063] The obtained Fe-Co-Ni-Al alloy system is named Ni650 alloy.

[0064] Example 5

[0065] Compared with Example 4, the only difference is that the aging temperature is 700 °C.

[0066] The obtained Fe-Co-Ni-Al alloy system is named Ni700 alloy.

[0067] Example 6

[0068] Compared with Example 4, the only difference is that the aging temperature is 750 °C.

[0069] The obtained Fe-Co-Ni-Al alloy system is named Ni750 alloy.

[0070] Comparative Example 1

[0071] Compared with Example 1, the only difference is that for Fe a Co b Ni c Al d Ti e Nb f Ta g B h a, b, c, d, e, f, g are 46.5%, 14%, 28%, 11.5%, 0%, 0%, 0%, 0% respectively.

[0072] The obtained Fe-Co-Ni-Al alloy system is named Alloy 650.

[0073] Comparative Example 2

[0074] Compared with Comparative Example 1, the only difference is that the aging temperature is 700 °C.

[0075] The obtained Fe-Co-Ni-Al alloy system is named Alloy 700.

[0076] Comparative Example 3

[0077] Compared with Comparative Example 1, the only difference is that the aging temperature is 750 °C.

[0078] The obtained Fe-Co-Ni-Al alloy system is named Alloy 750.

[0079] Comparative Example 4

[0080] Compared with Example 1, the only difference is that the aging annealing treatment in step (5) is not carried out.

[0081] The obtained Fe-Co-Ni-Al alloy system is named Co recrystallized alloy.

[0082] Comparative Example 5

[0083] Compared with Example 4, the only difference is that the aging annealing treatment in step (5) is not carried out.

[0084] The obtained Fe-Co-Ni-Al alloy system is named Ni recrystallized alloy.

[0085] In Examples 4 to 6, the alloy compositions are shown in Table 1:

[0086]

[0087] In Comparative Examples 1 to 3, the alloy compositions are shown in Table 2:

[0088]

[0089] Characterization of the materials prepared in Examples 1 to 6 and Comparative Examples 1 to 5:

[0090] For the high-entropy alloys with extremely high work hardening ability prepared in Examples 1 to 6 according to the described preparation method, room-temperature quasi-static tensile tests, synchrotron radiation diffraction tests (HEXRD), and scanning electron microscope imaging tests (SEM) were carried out. The test results are shown in Figures 1 to 6 .

[0091] (1) Room-temperature quasi-static tensile properties of Fe-Co-Ni-Al high-entropy alloys:

[0092] As shown Figure 1 in (a) and (b) below, the Co650 alloy obtained after melting, casting, cold rolling, recrystallization annealing, and aging treatment at 650 °C with Fe 43.96 Co 14 Ni 28 Al 11.5 Ti 0.83 Nb 0.83 Ta 0.83 B 0.05 as components has a work-hardening rate of up to 20 GPa at the plastic deformation stage, a yield strength of 1.2 GPa, a tensile strength of 2.4 GPa, and a uniform elongation of 19%, indicating that this example has extremely excellent comprehensive mechanical properties.

[0093] As shown Figure 2 in (a) and (b) below, the Co700 alloy obtained after adjusting the aging temperature to 700 °C has a further increased work-hardening rate at the plastic deformation stage, up to 25 GPa, a yield strength of 1 GPa, a tensile strength of 2.4 GPa, and a uniform elongation of 15%, indicating that this example also has excellent comprehensive mechanical properties and the increase in the aging temperature makes an important contribution to the improvement of the work-hardening rate.

[0094] As shown Figure 3 in (a) and (b) below, it shows that after adjusting the aging temperature to 750 °C, the work-hardening rate of the alloy reaches up to 22 GPa at the plastic deformation stage, the yield strength is 0.97 GPa, the tensile strength is 2.05 GPa, and the uniform elongation is 15%, indicating that as the aging temperature further increases, the strength and work-hardening ability of the alloy gradually decrease, but it still maintains good plasticity.

[0095] With Fe 43.96 Co 17 Ni 25 Al 11.5 Ti 0.83 Nb 0.83 Ta 0.83 B 0.05 as alloy components, Examples 4 to 6 also have good mechanical properties. The mechanical properties of the Ni650 alloy obtained by aging at 650 °C are as shown Figure 4As shown in (a) and (b) of Figure [Figure number not provided], the work hardening rate of Ni650 alloy can reach up to 30 GPa in the plastic deformation stage, the yield strength is 0.7 GPa, the tensile strength also reaches 2.5 GPa, and the uniform elongation rate is 13%. This indicates that after the composition adjustment, the work hardening ability of the alloy has been significantly improved. On the premise of maintaining a yield strength of 0.7 GPa and a uniform elongation rate of 13%, the tensile strength has been increased to an extremely excellent 2.5 GPa. After raising the aging temperature to 700 °C, Ni700 alloy is obtained, and its mechanical properties are as Figure 5 shown in (a) and (b) of Figure [Figure number not provided]. As can be seen from Figure 5 this, the work hardening rate of Ni700 alloy reaches the maximum value among all the examples of the present invention in the plastic deformation stage, up to 40 GPa, the yield strength is 0.5 GPa, the tensile strength reaches 2.3 GPa, and the uniform elongation rate is 10%. Figure 6 As shown in (a) and (b) of Figure [Figure number not provided], after adjusting the aging temperature to 750 °C, the work hardening rate of Ni750 alloy reaches up to 24 GPa in the plastic deformation stage, the yield strength is 0.5 GPa, the tensile strength is 2 GPa, and the uniform elongation rate is 12%. This indicates that with the further increase of the aging temperature, the strength and work hardening ability of the alloy gradually decrease, but it still maintains good plasticity.

[0096] (2) Characterization of the crystal structure and microstructure of the Fe-Co-Ni-Al high-entropy alloy with good mechanical properties:

[0097] Figures 1 to 6 Parts (c) and (d) of Figure [Figure number not provided] respectively correspond to the results of synchrotron radiation diffraction tests and scanning electron microscope imaging tests on Co650, Co700, Co700, Ni650, Ni700, and Ni750 alloys. As shown in the figure, Figure 1 diffraction peaks corresponding to the face-centered cubic crystal structure, body-centered cubic crystal structure, L12 ordered phase crystal structure, and B2 ordered phase crystal structure can be seen in (c) of Figure [Figure number not provided]. At the same time, as can be seen from Figure 1 (d) of Figure [Figure number not provided], equiaxed grains with uniform size distribution, fine black precipitate particles embedded between the equiaxed grain boundaries, and large-size island-shaped black phases randomly distributed can be seen. Among them, the equiaxed grain size is about 3.3 - 5.4 μm, the size of the fine black precipitate phase is about 1 - 2 μm, and the size of the island-shaped black phase is about 3 - 10 μm. This characterization result proves that the microstructure of this alloy is indeed a duplex microstructure composed of a face-centered cubic crystal structure phase and a B2 ordered body-centered cubic crystal structure phase, and there are nano-scale L12 ordered face-centered cubic crystal structure precipitate phases. Among them, the nano-scale L12 precipitate phases cannot be observed due to the resolution limit of the scanning electron microscope, but their existence can be proved by the results of synchrotron radiation diffraction tests. As shown in Figures 2 to 6 It should be noted that the figure numbers in the original text seem to be incomplete. You may need to check and supplement them accurately according to the actual situation.As shown in parts (c) and (d) thereof, the characterization results of the microcrystalline structures corresponding to the other five example alloys are all similar to those of Example 1, proving that the microstructural compositions of all examples are biphasic microstructures composed of a face-centered cubic crystal structure phase and a B2 ordered body-centered cubic crystal structure phase, and there are nanoscale L12 ordered face-centered cubic crystal structure precipitation phases.

[0098] (3)Comparison of the performance characterizations of Examples 1 to 5

[0099] Figure 7 is the tensile property of this alloy system without Ti, Nb, Ta, and B elements. The specific component composition of the alloy is: Fe a Co b Ni c Al d Ti e Nb f Ta g B h , where a, b, c, d, e, f, g are 46.5%, 14%, 28%, 11.5%, 0%, 0%, 0%, 0% respectively. Figure 7 In (a) and (b), they correspond to the tensile property curve and strain hardening curve of the 650 alloy, (c) and (d) correspond to the tensile property curve and strain hardening curve of the 700 alloy, and (e) and (f) correspond to the tensile property curve and strain hardening curve of the 750 alloy. The work hardening rates of the three alloys are all significantly reduced, and the highest work hardening rate of the 650 alloy, which is the only one with a work hardening rate increase after yielding, is only 4000 MPa. Therefore, it can be seen that Ti, Nb, Ta, and B elements play an indispensable role in ensuring the excellent work hardening performance of the alloy. Among them, the functions of Ti, Nb, and Ta elements are to form and stabilize the L12 element, and a trace amount of B element can help stabilize the grain boundary and inhibit the formation of harmful phases at the grain boundary.

[0100] Figure 8 In (a) and (b), they are the tensile property curve and work hardening curve of the Co recrystallized alloy; (c) and (d) are the tensile property curve and work hardening curve of the Ni recrystallized alloy. As can be seen from the figure, the plasticity of the two recrystallized alloys without aging treatment has increased, but the work hardening rate has decreased significantly. The highest work hardening rate after alloy yielding is only 5000 MPa. It shows that the aging process step in the present invention is indispensable, and its role is to promote the formation of L12 precipitation phases.

Claims

1. A Fe-Co-Ni-Al high-entropy alloy with extremely high work hardening ability, characterized in that: Its general formula for elemental composition is Fe a Co b Ni c Al d Ti e Nb f Ta g B h , where a, b, c, d, e, f, g, and h all represent the atomic percentages corresponding to each element, 40% ≤ a ≤ 50%, 10% ≤ b ≤ 20%, 20% ≤ c ≤ 30%, 5% ≤ d ≤ 15%, 0.4% ≤ e ≤ 2%, 0.4% ≤ f ≤ 2%, 0.4% ≤ g ≤ 2%, 0.01% ≤ h ≤ 1%, and a + b + c + d + e + f + g + h = 100%; it has a duplex microstructure of a face-centered cubic crystal structure phase and a B2 ordered body-centered cubic crystal structure phase, and there are nano-scale L12 ordered face-centered cubic crystal structure precipitation phases.

2. The Fe-Co-Ni-Al high-entropy alloy with extremely high work hardening ability according to claim 1, characterized in that: In the general formula of the element composition, 40% ≤ a ≤ 45%, 13% ≤ b ≤ 20%, 23% ≤ c ≤ 30%, 8% ≤ d ≤ 13%, 0.5% ≤ e ≤ 1%, 0.5% ≤ f ≤ 1%, 0.5% ≤ g ≤ 1%, 0.01% ≤ h ≤ 0.1%.

3. The preparation method of a Fe-Co-Ni-Al high-entropy alloy with extremely high work hardening ability according to claim 1 or 2, characterized in that: It is obtained by successively cleaning, vacuum melting, vacuum casting, solution treatment, cold rolling, recrystallization annealing treatment and aging treatment of Fe, Co, Ni, Al, Ti, Nb, Ta and B raw materials.

4. The preparation method of an Fe-Co-Ni-Al high-entropy alloy with extremely high work hardening ability according to claim 3, characterized in that: The cleaning process includes a step of grinding to remove the surface oxide layer and a step of ultrasonic washing.

5. The preparation method of an Fe-Co-Ni-Al high-entropy alloy with extremely high work hardening ability according to claim 3, characterized in that: The process of the vacuum melting: Each element is sorted according to its melting point and element type for the placement order. First, B is placed at the bottom, and the subsequent placement order of the remaining elements is Al, Ti, Ni, Co, Fe, Nb and Ta in sequence.

6. The preparation method of an Fe-Co-Ni-Al-based high-entropy alloy with extremely high work hardening ability according to claim 3 or 5, characterized in that: The conditions of the vacuum melting are: the arc current is 300 - 400 A, the melting time for each time is 80 - 120 s. After each melting, the ingot is turned over and cooled, and then the next melting is carried out. The melting is repeated at least eight times to make the alloy fully melted and uniform.

7. The preparation method of a Fe-Co-Ni-Al high-entropy alloy with extremely high work hardening ability according to claim 3, characterized in that: The conditions of the solution treatment are: at a temperature of 1200 - 1300 °C, after holding for 2 - 24 hours, water quenching is carried out.

8. The preparation method of an Fe-Co-Ni-Al high-entropy alloy with extremely high work hardening ability according to claim 3, characterized in that: The conditions of the cold rolling are: the ambient temperature is 25 - 30 °C, the reduction per pass is 3 - 5 mm, and the total rolling reduction is 80% - 90%.

9. The preparation method of an Fe-Co-Ni-Al high-entropy alloy with extremely high work hardening ability according to claim 3, characterized in that: The conditions of the recrystallization annealing treatment are: the temperature is 1000 - 1200 °C, and the time is 5 - 15 minutes.

10. The preparation method of an Fe-Co-Ni-Al high-entropy alloy with extremely high work hardening ability according to claim 3, characterized in that: The conditions of the aging treatment are: the temperature is 600 - 800 °C, and the time is 1 - 3 hours.

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