A Fe-Mn-Al-Ni based superelastic alloy and its preparation method

Abstract

This invention provides a Fe-Mn-Al-Ni based superelastic alloy and its preparation method, belonging to the technical field of iron-based superelastic alloys. By adding Co, this invention can form a high-density γ phase. After reheating back to the α single-phase region, a subgrain structure with a wide orientation distribution, high average orientation difference, and high number density is formed within the α phase grains. This leads to the aberrant growth of individual subgrains. The aberrantly grown subgrains completely consume other subgrains in the parent grain, creating a large free energy gradient between the parent grain and surrounding normal grains with subgrains. This provides a greater driving force for the growth of aberrant grains, significantly shortening the growth process. By adding Mo, the width of the low-density subgrain region can be increased, extending the time for grain boundaries to cross the low-density subgrain region, thereby reducing the number of aberrant grains and increasing the final grain size.

Description

Technical Field

[0001] This invention belongs to the field of iron-based superelastic alloy technology, specifically relating to a Fe-Mn-Al-Ni based superelastic alloy and its preparation method. Background Technology

[0002] Hyperelastic alloys, as intelligent and functional materials, have wide commercial applications in aerospace, biomedicine, civil engineering, construction, and vibration reduction. Ni-Ti based hyperelastic alloys are currently the most commercially mature type of hyperelastic alloy. However, the high cost and poor cold working properties of this system, with Ni and Ti accounting for up to 50 at.%, limit its further development.

[0003] Iron-based superelastic alloys are a type of superelastic alloy with iron as the matrix. Their superelastic effect manifests as a change in shape when stress is applied above the austenitic transformation completion temperature, followed by shape recovery after stress relief. They are inexpensive and possess excellent cold-working properties. Among them, Fe-Mn-Al-Ni superelastic alloys exhibit extremely low martensitic phase transformation entropy, with the critical martensitic phase transformation stress remaining almost unchanged within a temperature range of approximately 77–513 K, and also possessing a wide operating temperature window. These unique properties make Fe-Mn-Al-Ni superelastic alloys a powerful alternative to Ni-Ti-based superelastic alloys.

[0004] The reason why hyperelastic alloys can achieve shape recovery after stress loading and unloading is due to the thermoelastic martensitic phase transformation: with increasing stress, thermoelastic martensite laths nucleate and expand; with stress unloading, the thermoelastic martensite laths shrink and disappear. Studies on Fe-Mn-Al-Ni hyperelastic alloys show that grain boundaries inhibit the phase transformation by hindering the expansion of thermoelastic martensite laths, thus adversely affecting hyperelasticity and acting as the biggest obstacle to martensite lath growth and expansion. Therefore, the maximum recoverable strain of polycrystalline (average of more than two grains in the width / thickness direction of the sample) Fe-Mn-Al-Ni hyperelastic alloys is extremely low, only 2%~3%, which is far from meeting the requirements of industrial applications.

[0005] Reducing the number of grain boundaries (especially triplet grain boundaries) can effectively promote thermoelastic martensitic phase transformation. When the number of triplet grain boundaries is zero, i.e., when an oligocrystalline (only one grain in the width / thickness direction of the sample) or single-crystal structure is obtained, the alloy exhibits recoverable strain of over 7% and excellent superelasticity. However, normal grain growth can hardly completely eliminate triplet grain boundaries. Abnormal grain growth is a phenomenon in which an individual grain devours surrounding grains and grows rapidly. When an abnormal grain expands to completely cover the cross-section of the sample, an oligocrystalline structure can be obtained. Existing technology adds Co to Fe-Mn-Al-Ni superelastic alloys to form a high-density γ phase. After reheating back to the α single-phase region, a subcrystalline structure with a wide orientation distribution, high average orientation difference, and high number density is formed within the α phase grains. The wide orientation distribution of subgrains promotes the occurrence of abnormal grain growth. However, due to the increased number density of subgrains, the normal grain growth rate also increases accordingly. More importantly, the number of abnormal grain growth is uncontrollable, making it difficult to stably obtain an oligocrystalline structure, and thus difficult to achieve ideal superelasticity. Summary of the Invention

[0006] The purpose of this invention is to provide a Fe-Mn-Al-Ni based superelastic alloy and its preparation method. The Fe-Mn-Al-Ni based superelastic alloy provided by this invention can inhibit normal grain growth and reduce the number of abnormal grain growths, thereby improving the alloy's superelasticity.

[0007] To achieve the above-mentioned objectives, the present invention provides the following technical solution:

[0008] This invention provides a Fe-Mn-Al-Ni based superelastic alloy, comprising, by atomic percentage: Mn 25~40%, Al 10~20%, Ni 1~17%, Co 1~7.5%, Mo 0.1~1%, and the balance Fe.

[0009] Preferably, the Fe-Mn-Al-Ni based superelastic alloy comprises, by atomic percentage: Mn 30~35%, Al 13~17%, Ni 5~15%, Co 1.5~7.5%, Mo 0.2~1%, and the balance Fe.

[0010] Preferably, the Fe-Mn-Al-Ni based superelastic compound comprises, by atomic percentage: Mn 34%, Al 15%, Ni 7.5%, Co 2~7.5%, Mo 0.3~1%, and the balance Fe.

[0011] Preferably, the Fe-Mn-Al-Ni based superelastic compound comprises, by atomic percentage: Mn 34%, Al 15%, Ni 7.5%, Co 5~7.5%, Mo 0.4~1%, and the balance Fe.

[0012] This invention also provides a method for preparing the Fe-Mn-Al-Ni based superelastic alloy described in the above technical solution, comprising the following steps:

[0013] The Fe-Mn-Al-Ni-Co-Mo alloy castings were homogenized and then hot-rolled to obtain billets.

[0014] The billet is subjected to cyclic heat treatment or directional recrystallization to obtain Fe-Mn-Al-Ni based superelastic alloy.

[0015] Preferably, the cyclic heat treatment includes the following steps:

[0016] The billet is first heated to a first temperature, then cooled to a second temperature, and then heated to a third temperature.

[0017] After the second temperature is raised to the third temperature, it is cooled to room temperature, or the cooling and second temperature-raising process is repeated and then cooled to room temperature.

[0018] The first temperature is 1200~1300℃; the second temperature is not higher than 1180℃; and the third temperature is 1190~1200℃.

[0019] Preferably, the parameters for directional recrystallization include: a hot zone temperature of 1180~1300℃, a hot zone width of 1~100mm, a temperature gradient of 10~500℃ / mm, and a pulling speed of 0.1~1000μm / s.

[0020] Preferably, the homogenization treatment temperature is 1200~1300℃, and the holding time is 1~10h.

[0021] Preferably, the initial rolling temperature of the hot rolling is not lower than 1200℃, and the final rolling temperature of the hot rolling is not lower than 1100℃.

[0022] Preferably, the deformation amount of the hot rolling is not less than 50%.

[0023] This invention provides a Fe-Mn-Al-Ni based superelastic alloy, comprising, by atomic percentage: Mn 25~40%, Al 10~20%, Ni 1~17%, Co 1~7.5%, Mo 0.1~1%, and the balance Fe. Fe-Mn-Al-Ni based superelastic alloys are alloys with a high-temperature α single-phase region and a mid-temperature α+γ two-phase region. When the temperature is lowered from the α single-phase region to the α+γ two-phase region, held for a certain time, and then heated back to the α single-phase region, subcrystalline structures remain in the α phase. These subcrystalline structures are unevenly distributed in the grains, inducing abnormal grain growth and providing a driving force for it. This invention, by adding Co, lowers the formation energy of the γ phase itself without significantly affecting the α→γ transformation temperature, and can form a high-density γ phase. After heating back to the α single-phase region, subcrystalline structures with broad orientation distribution, high average orientation difference, and high number density are formed within the α phase grains. The broad orientation distribution of subcrystalline structures first leads to the abnormal growth of individual subcrystalline structures. The abnormally grown subcrystalline structures completely consume other subcrystalline structures in the parent grain, creating a huge free energy gradient between the parent grain and the surrounding normal grains with subcrystalline structures. Therefore, the grain boundaries of the parent grain migrate in the direction of decreasing free energy, making the parent grain always selectively grown. The growth of anomalous grains is induced by the high average orientation difference and number density of the subgrain structure, which further expands the free energy per unit volume in normal grains, providing a greater driving force for anomalous grain growth and significantly shortening the growth process. The addition of Mo increases the width of the subgrain low-density region, prolonging the time for grain boundaries to cross this region and extending the time for anomalous grain growth induced by it. This reduces the number of anomalous grains and increases the final grain size. Furthermore, the addition of Mo does not affect the anomalous growth of subgrains. Compared to existing Fe-Mn-Al-Ni superelastic alloys, the Fe-Mn-Al-Ni-based superelastic alloy of this invention synergistically improves grain growth efficiency and growth capacity. By limiting the amount of Co, grain growth efficiency and growth capacity are ensured. By limiting the amount of Mo, the width of the subgrain low-density region is increased while avoiding affecting the α→γ transition temperature. The results of the embodiments show that the grain boundary migration rate of the Fe-Mn-Al-Ni-based superelastic alloy provided by this invention is 21.1 × 10⁻⁶. -6 m / s, with a maximum recoverable strain of 7.2%. Attached Figure Description

[0024] Figure 1 These are the EBSD diagrams of alloys 1-3 of the comparative examples of this invention;

[0025] Figure 2 The metallographic and EBSD diagrams of alloys 4-5 of the present invention are shown below.

[0026] Figure 3 Metallographic images of alloys 6-8 of the comparative examples of this invention;

[0027] Figure 4 Metallographic images of alloys from Comparative Examples 9-10 and Example 2 of this invention;

[0028] Figure 5 This is a graph showing the change in grain boundary migration of different alloys of the present invention with reheating time;

[0029] Figure 6 The stress-strain curves are for the alloys of Comparative Example 15 and Example 5 of this invention. Detailed Implementation

[0030] All raw materials used in this invention are not particularly limited in their source; they can be purchased from the market or prepared using conventional methods known to those skilled in the art.

[0031] There are no particular restrictions on the purity of any of the raw materials used in this invention, but high-purity raw materials are preferred.

[0032] This invention provides a Fe-Mn-Al-Ni based superelastic alloy, comprising, by atomic percentage: Mn 25~40%, Al 10~20%, Ni 1~17%, Co 1~7.5%, Mo 0.1~1%, and the balance Fe.

[0033] The Fe-Mn-Al-Ni based superelastic alloy provided by this invention, by atomic percentage, comprises 25-40% Mn, preferably 30-35%. As one embodiment of this invention, the atomic percentage of Mn in the Fe-Mn-Al-Ni based superelastic alloy can be 25%, 30%, 34%, 35%, or 40%. When the atomic percentage of Mn is within the above range, it can synergize with other elements to ensure the superelasticity of the alloy.

[0034] The Fe-Mn-Al-Ni based superelastic alloy provided by this invention, by atomic percentage, comprises 10-20% Al, preferably 13-17%. As one embodiment of this invention, the atomic percentage of Al in the Fe-Mn-Al-Ni based superelastic alloy can be 10%, 12%, 15%, 18%, or 20%. When the atomic percentage of Al is within the above range, it can synergize with other elements to ensure the superelasticity of the alloy.

[0035] The Fe-Mn-Al-Ni based superelastic alloy provided by this invention, by atomic percentage, comprises 1-17% Ni, preferably 5-15%. As one embodiment of this invention, the atomic percentage of Ni in the Fe-Mn-Al-Ni based superelastic alloy can be 2%, 5%, 7.5%, 10%, 15%, or 17%. When the atomic percentage of Ni is within the above range, it can synergize with other elements to ensure the superelasticity of the alloy.

[0036] The Fe-Mn-Al-Ni based superelastic alloy provided by this invention, by atomic percentage, comprises 1-7.5% Co, preferably 1.5-7.5%, more preferably 2-7.5%, and even more preferably 5-7.5%. As one embodiment of this invention, the atomic percentage of Co in the Fe-Mn-Al-Ni based superelastic alloy can be 1%, 2%, 3%, 4%, 5%, 6%, or 7.5%. The atomic radius of Co is close to that of Fe, Mn, Al, and Ni, and Co has extremely high solid solubility in α-Fe, thus not affecting the brittleness of the alloy. Adding Co to the Fe-Mn-Al-Ni based alloy can reduce the formation energy of BCC+FCC cells. ΔE Lowering the formation energy of the γ phase while having almost no impact on the α→γ transformation temperature allows for the formation of a high-density γ phase. Upon reheating back to the α single-phase region, a subcrystalline structure with broad orientation distribution, high average orientation difference, and high number density is formed within the α phase grains. The broad orientation distribution of subcrystalline grains initially leads to the aberrant growth of individual subcrystalline grains. The aberrantly grown subcrystalline grains completely consume other subcrystalline grains in the parent grain, creating a large free energy gradient between the parent grain and the surrounding normal grains with subcrystalline grains. Therefore, the grain boundaries of the parent grain migrate in the direction of decreasing free energy, ensuring that the parent grain is always selected as an aberrant grain and undergoes growth. The high average orientation difference and high number density of the subcrystalline structure further expand the unit volume free energy in the normal grains filled with subcrystalline grains, providing a greater driving force for the growth of aberrant grains and significantly shortening the growth process of aberrant grains. With the atomic percentage of Co within the above range, the growth efficiency and growth capacity of the grains can be guaranteed, while avoiding excessive content that would increase costs.

[0037] The Fe-Mn-Al-Ni based superelastic alloy provided by this invention, by atomic percentage, comprises 0.1-1% Mo, preferably 0.2-1%, more preferably 0.3-1%, and even more preferably 0.4-1%. As one embodiment of this invention, the atomic percentage of Mo in the Fe-Mn-Al-Ni based superelastic alloy can be 0.1%, 0.2%, 0.3%, 0.4%, 0.5%, 0.8%, or 1%. Adding Mo to the Fe-Mn-Al-Ni alloy can increase the width of the subgrain low-density region, prolonging the time for grain boundaries to cross the subgrain low-density region, thus extending the time for aberrant grain growth induced by the subgrain low-density region to occur, thereby reducing the number of aberrant grains and increasing the final grain size. Furthermore, the addition of Mo does not affect the aberrant growth of subgrains. With the atomic percentage of Mo within the above range, the width of the subgrain low-density region can be increased while avoiding affecting the α→γ transformation temperature (a decrease in the α→γ transformation temperature will affect the processing temperature range of the alloy).

[0038] The Fe-Mn-Al-Ni based superelastic alloy provided by this invention also includes the balance Fe. As the matrix of the alloy, Fe can work synergistically with other elements to give the alloy superelasticity and reduce the cost of the alloy.

[0039] Fe-Mn-Al-Ni based superelastic alloys are alloys with a high-temperature α single-phase region and a mid-temperature α+γ two-phase region. When the temperature is lowered from the α single-phase region to the α+γ two-phase region, held for a certain time, and then heated back to the α single-phase region, subcrystalline structures remain in the α phase. These subcrystalline structures are unevenly distributed in the grains, inducing abnormal grain growth and providing a driving force for it. This invention, by adding Co, lowers the formation energy of the γ phase itself without significantly affecting the α→γ transformation temperature, and can form a high-density γ phase. After heating back to the α single-phase region, subcrystalline structures with broad orientation distribution, high average orientation difference, and high number density are formed within the α phase grains. The broad orientation distribution of subcrystalline structures first leads to the abnormal growth of individual subcrystalline structures. The abnormally grown subcrystalline structures completely consume other subcrystalline structures in the parent grain, creating a huge free energy gradient between the parent grain and the surrounding normal grains with subcrystalline structures. Therefore, the grain boundaries of the parent grain migrate in the direction of decreasing free energy, making the parent grain always selected as the... Abnormal grains grow; the high average orientation difference and number density of the subgrain structure further expand the unit volume free energy in the normal grains filled with subgrains, providing a greater driving force for grain growth and significantly shortening the process of abnormal grain growth; by adding Mo, the width of the subgrain low-density region can be increased, extending the time for grain boundaries to cross the subgrain low-density region, thus prolonging the time for abnormal grain growth induced by the subgrain low-density region to occur, thereby reducing the number of abnormal grains and increasing the final grain size, and the addition of Mo does not affect the abnormal growth of subgrains; compared with existing Fe-Mn-Al-Ni superelastic alloys, the Fe-Mn-Al-Ni based superelastic alloy of this invention synergistically improves the grain growth efficiency and growth capacity; by limiting the amount of Co, the grain growth efficiency and growth capacity are guaranteed; by limiting the amount of Mo, the width of the subgrain low-density region is increased while avoiding affecting the α→γ transformation temperature.

[0040] This invention also provides a method for preparing the Fe-Mn-Al-Ni based superelastic alloy described in the above technical solution, comprising the following steps:

[0041] The Fe-Mn-Al-Ni-Co-Mo alloy castings were homogenized and then hot-rolled to obtain billets.

[0042] The billet is subjected to cyclic heat treatment or directional recrystallization to obtain Fe-Mn-Al-Ni based superelastic alloy.

[0043] This invention involves homogenizing Fe-Mn-Al-Ni-Co-Mo alloy castings and then hot rolling them to obtain billets.

[0044] As one embodiment of the present invention, the Fe-Mn-Al-Ni-Co-Mo alloy casting can be obtained by vacuum melting; the vacuum degree of the vacuum melting can be below 0.1 Pa; the vacuum melting can be carried out in a vacuum non-consumable arc melting furnace.

[0045] In one embodiment of the present invention, the raw materials for the Fe-Mn-Al-Ni-Co-Mo alloy are metallic iron, metallic manganese, metallic aluminum, metallic nickel, metallic cobalt, and metallic molybdenum. The vacuum melting process can be as follows: first, a layer of pure Fe is spread evenly in the melting crucible; then, volatile Mn is added, followed by low-melting-point Al; then Ni, Co, and Mo are added; finally, the remaining Fe is added. Argon gas is introduced as both the arc-igniting medium and the protective gas. During the melting process, magnetic stirring is activated to ensure thorough mixing of the alloy liquid. The melting process is repeated 3-4 times to ensure the uniformity of the casting and prevent compositional segregation.

[0046] In this invention, the homogenization treatment temperature is preferably 1200~1300℃, more preferably 1200~1250℃; the homogenization treatment holding time is preferably 1~10h, more preferably 1~5h. Homogenization treatment parameters within the above ranges can ensure uniform distribution of elements and reduce component segregation.

[0047] As one embodiment of the present invention, the homogenization process can be carried out in air, vacuum or a protective atmosphere.

[0048] As one embodiment of the present invention, the cooling method for the homogenization process can be air cooling or water cooling.

[0049] In one embodiment of the present invention, the casting can be heated to 1200~1300℃ and held for 1~10 hours before hot rolling.

[0050] In this invention, the initial rolling temperature is preferably not lower than 1200°C, and the final rolling temperature is preferably not lower than 1100°C. Controlling the hot rolling temperature within the above range is beneficial for promoting element diffusion, reducing as-cast dendrites, and further improving uniformity.

[0051] In this invention, the deformation amount of the hot rolling is preferably not less than 50%. Controlling the deformation amount of the hot rolling within the above range is beneficial to promoting element diffusion, reducing cast dendrites, and further improving uniformity.

[0052] As one embodiment of the present invention, the hot-rolled material can be air-cooled or water-cooled to room temperature.

[0053] After obtaining the billet, the present invention performs cyclic heat treatment or directional recrystallization on the billet to obtain Fe-Mn-Al-Ni based superelastic alloy.

[0054] In one embodiment of the present invention, the blank can be cut first. The present invention does not have special requirements on the cutting dimensions; cutting can be performed according to needs. In embodiments of the present invention, the cutting dimensions are 200mm × 25mm and 4000mm × 100mm.

[0055] In one technical solution of the present invention, the billet is subjected to cyclic heat treatment; the cyclic heat treatment preferably includes the following steps:

[0056] The billet is first heated to a first temperature, then cooled to a second temperature, and then heated to a third temperature.

[0057] After the second temperature rises to the third temperature, the temperature is cooled to room temperature, or the cooling and second temperature rise processes are repeated and then cooled to room temperature.

[0058] In this invention, the first temperature is preferably 1200~1300℃, more preferably 1200~1250℃; the second temperature is preferably not higher than 1180℃, more preferably 900~1000℃; and the third temperature is preferably 1190~1200℃. Controlling the temperature at different stages within the above ranges is beneficial for controlling the number of abnormal grain growths and further improving the superelasticity of the alloy.

[0059] As one embodiment of the present invention, after the temperature reaches the first temperature, the second temperature and the third temperature, heat preservation can be performed; the present invention does not have a specific limitation on the heat preservation time, and the heat preservation time can be adjusted according to actual needs.

[0060] In one embodiment of the present invention, the cyclic heat treatment step can be as follows: heating the billet to 1200℃, holding it at that temperature for 60 min, cooling it to 900℃, holding it at that temperature for 15 min, then heating it to 1190℃ or 1200℃, without holding it at that temperature, or holding it at that temperature for 3.3 min or 15 min, and then cooling it to room temperature. The α-γ transformation temperature of the Fe-Mn-Al-Ni-Co-Mo superelastic alloy provided by the present invention is 1180℃. Heating the billet to 1200~1300℃ first can fully transform the alloy into the α phase. Cooling it to below 1180℃ can form an α+γ dual phase. Heating it to above 1190℃ can transform the γ phase into the α phase, leaving a subcrystalline structure in the α phase. Since the α→γ transformation has already begun at 1180℃, the remaining subcrystalline structure in the α phase can induce abnormal grain growth during the continued heating process.

[0061] In one embodiment of the present invention, the heating rate of each heating in the cyclic heat treatment can be 0.1~100℃ / min, or 1~5℃ / min, specifically 1℃ / min, 3℃ / min, 5℃ / min, 7℃ / min or 9℃ / min.

[0062] In one embodiment of the present invention, the cooling rate in the cyclic heat treatment can be 0.1~100℃ / min, or 1~10℃ / min, specifically 3℃ / min.

[0063] As one embodiment of the present invention, the cooling to room temperature during the cyclic heat treatment can be achieved by air cooling or water quenching.

[0064] In another technical solution of the present invention, the billet is subjected to directional recrystallization; the parameters of the directional recrystallization preferably include: a hot zone temperature of 1180~1300℃, a hot zone width of 1~100mm, a temperature gradient of 10~500℃ / mm, and a pulling rate of 0.1~1000μm / s. As an embodiment of the present invention, the parameters of the directional recrystallization can be: a hot zone temperature of 1250℃, a hot zone width of 9mm, a temperature gradient of 200℃ / mm, and a pulling rate of 7μm / s. The parameters of the directional recrystallization are within the above ranges, which can ensure sufficient phase transformation.

[0065] In one embodiment of the present invention, the directional recrystallization can be carried out under vacuum conditions, wherein the vacuum degree of the vacuum conditions can be 10. -4 Pa.

[0066] In one embodiment of the present invention, the directional recrystallization can be performed by connecting a servo motor in a directional annealing furnace.

[0067] As one embodiment of the present invention, the cooling method after directional recrystallization can be furnace cooling, air cooling, water cooling or oil cooling.

[0068] The preparation methods provided by this invention are all common processes, which make it easy to control product quality and are conducive to industrial production.

[0069] The technical solutions of this invention will be clearly and completely described below with reference to the embodiments thereof. Obviously, the described embodiments are only a part of the embodiments of this invention, and not all of them. All other embodiments obtained by those skilled in the art based on the embodiments of this invention without creative effort are within the scope of protection of this invention.

[0070] Example 1

[0071] A Fe-Mn-Al-Ni based superelastic alloy, by atomic percentage, is: Fe 41%, Mn 34%, Al 15%, Ni 7.5%, Co 1.5%, Mo 1%.

[0072] The preparation method is as follows:

[0073] Weigh out the pure metal raw materials according to the specified proportions;

[0074] First, a layer of pure Fe is spread on the melting crucible of the vacuum arc melting furnace, then Mn, Al, Ni, Co, and Mo are added in sequence, and finally the remaining Fe is added. Argon gas is introduced, and the vacuum is evacuated to 0.1 Pa to start melting. Magnetic stirring is turned on during the melting process. The casting is obtained after melting 4 times.

[0075] The casting was heated to 1200℃ in an argon atmosphere and held for 5 hours for homogenization, then air-cooled to room temperature. The homogenized casting was then heated to 1200℃ and held for 1 hour before hot rolling. The initial rolling temperature was 1200℃, and the final rolling temperature was 1100℃. The casting thickness was rolled from 150mm to 4mm, and then air-cooled to room temperature. The rolled samples were cut into 200mm×25mm×4mm and 4000mm×100mm×4mm billets, respectively, and installed in a directional annealing furnace connected to a servo motor pull rod. The furnace cover was closed, and a vacuum was applied to 10... -4 Pa, the hot zone was heated to 1250℃ using an induction heating coil before directional recrystallization began; the directional recrystallization parameters were: hot zone width 9mm, temperature gradient 200℃ / mm, and pulling speed 7μm / s; for smaller samples of 200mm×25mm×4mm, the directional recrystallization distance was 7.5cm, and for larger samples of 4000mm×100mm×4mm, the directional recrystallization distance was 20cm; after directional recrystallization, the sample was removed and air-cooled to obtain Fe-Mn-Al-Ni based superelastic alloy.

[0076] Example 2

[0077] A Fe-Mn-Al-Ni based superelastic alloy, by atomic percentage, comprises: Fe 35%, Mn 34%, Al 15%, Ni 7.5%, Co 7.5%, and Mo 1%.

[0078] The preparation method is as follows:

[0079] Weigh out the pure metal raw materials according to the specified proportions;

[0080] First, a layer of pure Fe is spread on the melting crucible of the vacuum arc melting furnace, then Mn, Al, Ni, Co, and Mo are added in sequence, and finally the remaining Fe is added. Argon gas is introduced, and the vacuum is evacuated to 0.1 Pa to start melting. Magnetic stirring is turned on during the melting process. The casting is obtained after melting 4 times.

[0081] The casting was heated to 1200℃ in an argon atmosphere and held for 5 hours for homogenization treatment, then air-cooled to room temperature. The homogenized casting was then heated to 1200℃ and held for 1 hour before hot rolling. The initial rolling temperature was 1200℃, and the final rolling temperature was 1100℃. The casting thickness was rolled from 20mm to 3mm, and then air-cooled to room temperature. The rolled billet underwent cyclic heat treatment: a first heating rate of 3℃ / min to 1200℃ and holding for 1 hour, followed by a cooling rate of 3℃ / min to 900℃ and holding for 15 minutes, and then a second heating rate of 3℃ / min to 1190℃ followed by water quenching to obtain the Fe-Mn-Al-Ni-Co-Mo superelastic alloy.

[0082] Example 3

[0083] A Fe-Mn-Al-Ni based superelastic alloy, by atomic percentage, comprises: Fe 35%, Mn 34%, Al 15%, Ni 7.5%, Co 7.5%, and Mo 1%.

[0084] The preparation method is the same as in Example 2, except that the reheating temperature during the cyclic heat treatment is 1200℃.

[0085] Example 4

[0086] A Fe-Mn-Al-Ni based elastic alloy, by atomic percentage, comprises: Fe 35%, Mn 34%, Al 15%, Ni 7.5%, Co 7.5%, and Mo 1%.

[0087] The preparation method is the same as in Example 2, except that the reheating temperature during the cyclic heat treatment is 1200℃ and held for 3.3 min.

[0088] Example 5

[0089] A Fe-Mn-Al-Ni based superelastic alloy, by atomic percentage, comprises: Fe 35%, Mn 34%, Al 15%, Ni 7.5%, Co 7.5%, and Mo 1%.

[0090] The preparation method is the same as in Example 2, except that the reheating temperature during the cyclic heat treatment is 1200℃ and held for 15 minutes.

[0091] Example 6

[0092] A Fe-Mn-Al-Ni based superelastic alloy, by atomic percentage, comprises: Fe 35%, Mn 34%, Al 15%, Ni 7.5%, Co 7.5%, and Mo 1%.

[0093] The preparation method is the same as in Example 2, except that the reheating temperature during the cyclic heat treatment is 1200℃ and held for 15 minutes, and the cyclic heat treatment cooling-reheating process is repeated 4 times.

[0094] Comparative Example 1

[0095] A Fe-Mn-Al-Ni-Co superelastic alloy, by atomic percentage, comprises: Fe 41.5%, Mn 34%, Al 15%, Ni 7.5%, and Co 2%.

[0096] The preparation method is the same as in Example 2, except that during the cyclic heat treatment, the temperature was lowered to 900°C and held for 15 minutes without reheating, and the sample was directly quenched in water.

[0097] Comparative Example 2

[0098] A Fe-Mn-Al-Ni-Co superelastic alloy, by atomic percentage, has the following composition: Fe 36%, Mn 34%, Al 15%, Ni 7.5%, and Co 7.5%.

[0099] The preparation method is the same as that of Comparative Example 1.

[0100] Comparative Example 3

[0101] A Fe-Mn-Al-Ni-Co-Mo superelastic alloy, by atomic percentage, comprises: Fe 35%, Mn 34%, Al 15%, Ni 7.5%, Co 7.5%, and Mo 1%.

[0102] The preparation method is the same as that of Comparative Example 1.

[0103] Comparative Example 4

[0104] A Fe-Mn-Al-Ni-Co superelastic alloy, by atomic percentage, comprises: Fe 41.5%, Mn 34%, Al 15%, Ni 7.5%, and Co 2%.

[0105] The preparation method is the same as in Example 2, except that the reheating temperature during the cyclic heat treatment is 1180°C.

[0106] Comparative Example 5

[0107] A Fe-Mn-Al-Ni-Co superelastic alloy, by atomic percentage, has the following composition: Fe 36%, Mn 34%, Al 15%, Ni 7.5%, and Co 7.5%.

[0108] The preparation method is the same as in Example 2, except that the reheating temperature during the cyclic heat treatment is 1180℃ and held for 1 minute.

[0109] Comparative Example 6

[0110] A Fe-Mn-Al-Ni-Mo superelastic alloy, by atomic percentage, comprises: Fe 42.5%, Mn 34%, Al 15%, Ni 7.5%, and Mo 1%.

[0111] The preparation method is the same as in Example 2, except that the reheating temperature during the cyclic heat treatment is 1180℃ and held for 2 minutes.

[0112] Comparative Example 7

[0113] A Fe-Mn-Al-Ni-Co superelastic alloy, by atomic percentage, has the following composition: Fe 36%, Mn 34%, Al 15%, Ni 7.5%, and Co 7.5%.

[0114] The preparation method is the same as in Example 2, except that the reheating temperature during the cyclic heat treatment is 1180℃ and held for 2 minutes.

[0115] Comparative Example 8

[0116] A Fe-Mn-Al-Ni-Co-Mo superelastic alloy, by atomic percentage, comprises: Fe 35%, Mn 34%, Al 15%, Ni 7.5%, Co 7.5%, and Mo 1%.

[0117] The preparation method is the same as in Example 2, except that the reheating temperature during the cyclic heat treatment is 1180℃ and held for 2 minutes.

[0118] Comparative Example 9

[0119] A Fe-Mn-Al-Ni-Mo superelastic alloy, by atomic percentage, comprises: Fe 42.5%, Mn 34%, Al 15%, Ni 7.5%, and Mo 1%.

[0120] The preparation method is the same as in Example 2.

[0121] Comparative Example 10

[0122] A Fe-Mn-Al-Ni-Co superelastic alloy, by atomic percentage, has the following composition: Fe 36%, Mn 34%, Al 15%, Ni 7.5%, and Co 7.5%.

[0123] The preparation method is the same as in Example 2.

[0124] Comparative Example 11

[0125] A Fe-Mn-Al-Ni superelastic alloy, by atomic percentage, comprises: Fe 43.5%, Mn 34%, Al 15%, and Ni 7.5%.

[0126] The preparation method is the same as that of Comparative Example 4.

[0127] Comparative Example 12

[0128] A Fe-Mn-Al-Ni superelastic alloy, by atomic percentage, comprises: Fe 43.5%, Mn 34%, Al 15%, and Ni 7.5%.

[0129] The preparation method is the same as in Example 2.

[0130] Comparative Example 13

[0131] A Fe-Mn-Al-Ni superelastic alloy, by atomic percentage, comprises: Fe 43.5%, Mn 34%, Al 15%, and Ni 7.5%.

[0132] The preparation method is the same as in Example 3.

[0133] Comparative Example 14

[0134] A Fe-Mn-Al-Ni superelastic alloy, by atomic percentage, comprises: Fe 43.5%, Mn 34%, Al 15%, and Ni 7.5%.

[0135] The preparation method is the same as in Example 4.

[0136] Comparative Example 15

[0137] A Fe-Mn-Al-Ni superelastic alloy, by atomic percentage, comprises: Fe 43.5%, Mn 34%, Al 15%, and Ni 7.5%.

[0138] The preparation method is the same as in Example 5.

[0139] Comparative Example 16

[0140] A Fe-Mn-Al-Ni superelastic alloy, by atomic percentage, comprises: Fe 43.5%, Mn 34%, Al 15%, and Ni 7.5%.

[0141] The preparation method is the same as in Example 6.

[0142] Comparative Example 17

[0143] A Fe-Mn-Al-Ni-Co-Mo superelastic alloy, by atomic percentage, comprises: Fe 35%, Mn 34%, Al 15%, Ni 7.5%, Co 7.5%, and Mo 1%.

[0144] The preparation method is the same as in Example 2, except that the reheating temperature during the cyclic heat treatment is 1180°C.

[0145] Test Example 1

[0146] The Fe-Mn-Al-Ni based superelastic alloy samples obtained in Example 1 were mechanically polished and then etched with a 5wt% nitric acid alcohol solution to observe the grain boundaries. Typically, in directional recrystallization, due to the additional driving force provided by surface energy, grains in small-sized samples tend to grow easily with the movement of the heating zone, while grains in large-sized samples, due to lower surface energy, are less likely to grow with the movement of the heating zone. In the smaller 200mm × 25mm × 4mm sample, the maximum grain length reached 7.5cm, the same as the directional recrystallization distance; in the larger 4000mm × 100mm × 4mm sample, even with a directional recrystallization distance of 20cm, the maximum grain length still matched well, indicating that this invention is independent of size effects.

[0147] Test Example 2

[0148] The superelastic alloys obtained in Comparative Examples 1-3 were observed by electron backscatter diffraction (EBSD), and the images obtained are as follows. Figure 1 As shown, a is the phase composition diagram of alloy 1 (Comparative Example), b is the phase composition diagram of alloy 2 (Comparative Example), c is the phase composition diagram of alloy 3 (Comparative Example), a1 is the standard grain orientation deviation diagram of alloy 1 (Comparative Example), b1 is the standard grain orientation deviation diagram of alloy 2 (Comparative Example), and c1 is the standard grain orientation deviation diagram of alloy 3 (Comparative Example). In the phase composition diagrams, red represents the α phase and black represents the γ phase.

[0149] according to Figure 1 Using image analysis software, the area fraction of γ-precipitate in the 2Co alloy (Comparative Example 1) was approximately 80%, and the number density was approximately 4.21 × 10⁻⁶. -8 / m 2 The γ-precipitate area fraction of the 7.5Co alloy (Comparative Example 2) was also around 80%, but the number density increased to 7.45 × 10⁻⁶. -8 / m 2 The γ-precipitate area fraction of the 7.5Co1Mo alloy (Comparative Example 3) was also around 80%, and its number density was roughly the same as that of the 7.5Co alloy, at 8.12 × 10⁻⁶. -8 / m 2 This indicates that Co increases the number density of γ-precipitates, while Mo has little effect on the characteristics of γ-precipitates.

[0150] according to Figure 1Compared to the 2Co alloy, in the 7.5Co and 7.5Co1Mo alloys, as the number density of γ precipitate increases, the reference orientation deviation angle of grains in the α matrix shows significant fluctuations, i.e., the red area gradually increases. This indicates that subgrains have already formed in the α matrix at this point, and the number of subgrains with high orientation difference increases with the increase of Co element. The addition of Mo element has little effect on the change of orientation difference of subgrains in Co-containing alloys.

[0151] Test Example 3

[0152] The superelastic alloys obtained in Comparative Examples 4 and 5 were observed using optical metallographic microscopy and EBSD, and the images are shown below. Figure 2 As shown, a is the metallographic image of Comparative Example 4 alloy, b is the metallographic image of Comparative Example 5 alloy, c, c1, and c2 are the image quality image, EBSD image, and standard grain orientation deviation image of a special region of the α matrix in Comparative Example 5 alloy, d is a magnified view of c, e is the inverse pole image of the two purple grains in d, a1 is the EBSD image of the α matrix of Comparative Example 4 alloy, b1 is the EBSD image of the α matrix of Comparative Example 5 alloy, a2 is the subgrain boundary orientation difference distribution image of Comparative Example 4 alloy, and b2 is the subgrain boundary orientation difference distribution image of Comparative Example 5 alloy.

[0153] from Figure 2 In images a and b, black spiderweb-like subgrains are clearly visible. A small amount of incompletely dissolved black γ-precipitate is also present in the 7.5Co alloy (Comparative Example 5); this is attributed to Co being an FCC stabilizing element, which slightly increases the dissolution temperature of the γ-precipitate. Figure 2 The average size of the subgrains in the 2Co alloy (Comparative Example 4) was measured to be 32 μm in both a and b samples, while it decreased to 15 μm in the 7.5Co alloy. This indicates that the number density of subgrains increases with increasing Co content.

[0154] from Figure 2 As can be seen from a1 and b1, a distinct subcrystalline structure can be observed inside the grain.

[0155] from Figure 2 As can be seen from a2 and b2, the subgrain boundary orientation difference of the 2Co alloy ranges from 0° to 3.5°, with an average orientation difference of 1.03°. In the 7.5Co alloy, the subgrain boundary orientation difference range widens to 0°–5°, and the average orientation difference increases to 1.34°. Combined with comparative examples 1 and 2, this indicates that Co element expands the subgrain boundary orientation difference range by increasing the number density of γ-precipitates, and also increases the average orientation difference and density of the subgrains.

[0156] from Figure 2 As can be seen from c, subgrains with very large differences in size and orientation (size ~100μm, orientation difference ~6.24°) were detected in this region, indicating that the subgrains have grown abnormally. Figure 2 In diagram d, three grains are highlighted with dashed lines: "Growth Grain", "Grain 1", and "Grain 2". ",according to Figure 2 The inverse pole figures of the two purple grains in e indicate "grain 1" and "grain 2". "Originally a single grain, the 'growing grain' shown by the blue dashed line is engulfing the purple grain and growing, and there are no subgrains within this grain. This indicates that the subgrains within this grain are consumed due to the abnormal growth of the subgrains, after which the grain grows preferentially. This is because once the subgrains in the grain are consumed due to the abnormal growth of the subgrains, there are no subgrains in this grain, while the adjacent normal grain contains subgrains. Therefore, there will be a free energy gradient on both sides of the grain boundary between the grain without subgrains and the normal grain, caused by the difference in subgrain density. Under the action of this gradient, the grain boundary begins to migrate towards the side with higher subgrain energy to reduce the system energy, that is, towards the side of the normal grain. Therefore, the grain without subgrains grows; moreover, there is a free energy gradient on each grain boundary of this grain, which promotes the continuous growth of the grain. It should be noted that even though an abnormal grain has been detected at this moment, the size of the abnormal grain is not significantly different from that of the normal grain, which means that rapid abnormal growth has not yet begun. The period before this moment is the 'latency period' of abnormal grain growth."

[0157] This phenomenon was not detected in Comparative Example 4 because the subgrain orientation difference range was small, and the difference between subgrain orientation differences was small. Since the migration rate of subgrain boundaries is proportional to the subgrain boundary energy, and the subgrain boundary energy is also proportional to the subgrain boundary orientation difference, the small subgrain boundary orientation difference makes the difference in migration rate between subgrain boundaries not significant. However, in Comparative Example 5, the large subgrain orientation difference range leads to a huge difference in migration rate between subgrain boundaries. Subgrain boundaries with large orientation differences have a more obvious kinetic advantage than subgrain boundaries with small orientation differences, thus causing abnormal subgrain growth and subsequent abnormal grain growth.

[0158] Test Example 4

[0159] The metallographic structures of the superelastic alloys in Comparative Examples 6-8 were observed using a metallographic microscope, and the images are as follows. Figure 3 As shown, a is a metallographic micrograph of alloy 6 (comparative example), b is a metallographic micrograph of alloy 7 (comparative example), and c is a metallographic micrograph of alloy 8 (comparative example).

[0160] from Figure 3 As can be seen from a and b, the subgrain size of the 1Mo alloy (Comparative Example 6) (~54.1 μm) is larger than that of the 7.5Co alloy (Comparative Example 7) (~15 μm), but the subgrain distribution characteristics of the two alloys are the same; the subgrains of both alloys have a slightly higher number density near the grain center compared to near the grain boundaries; and Figure 3 c shows that, compared to 1Mo and 7.5Co alloys, although the 7.5Co1Mo alloy (Comparative Example 8) also has high-density subgrains at the grain center, the subgrain density around the grain boundaries is extremely low and almost unobservable; the width of the extremely low-density subgrain region of the 7.5Co1Mo alloy is over 100 μm, while the width of the low-density subgrain region of both 1Mo and 7.5Co alloys is less than 50 μm; this indicates that neither Mo nor Co elements alone can significantly affect the distribution of subgrains, but the combination of Mo and Co elements leads to an increase in the width of the low-density subgrain region.

[0161] Test Example 5

[0162] The metallographic structure of the superelastic alloys obtained in Comparative Examples 9-10 and Example 2 was observed using a metallographic microscope, and the images are shown below. Figure 4 As shown, a is a metallographic micrograph of alloy of Comparative Example 9, b is a metallographic micrograph of alloy of Comparative Example 10, b1 is a partial magnified view of b, and c is a metallographic micrograph of alloy of Example 2.

[0163] from Figure 4 It can be seen that the small grain size and rich subgrain structure of the 1Mo alloy (Comparative Example 9) indicate that the abnormal grain growth process has not yet begun, while the 7.5Co alloy (Comparative Example 10) and the 7.5Co1Mo alloy (Example 2) have several large grains (greater than 4 mm), and the large grains almost cover the entire sample, indicating that the abnormal grain growth process is nearing completion. This shows that Co advances the time of abnormal grain growth compared to Mo, and the addition of Mo to Co-containing alloys does not change this advance effect. Figure 4 As shown in Figure b1, a small anomalous grain (~1.5 mm) with a typical subgrain-low density region was detected in the 7.5Co alloy. This grain has a high-density subgrain in its lower half and almost no subgrain structure in its upper half. Current research suggests that the subgrain-low density region structure near grain boundaries (such as...) Figure 3 As shown, this leads to anomalous grain growth. The width of the subgrain low-density region on both sides of each grain boundary is inconsistent. Grain boundaries with a favorable width in the subgrain low-density region (narrow subgrain low-density region) can preferentially contact the high-density subgrains at the grain center, thus having a kinetic advantage and causing anomalous grain growth. However, under this mechanism, since the width of the subgrain low-density region near the grain boundary is random, it is impossible to control which grain boundary will migrate rapidly. This means that even if a grain boundary migrates rapidly and one of the grains with that boundary is selected as an anomalous grain, the other grain boundaries of that grain may not have a kinetic advantage. Therefore, the migration mode of the other grain boundaries of that grain may be a slow, normal grain growth mode, resulting in low anomalous grain growth efficiency and small size under this mechanism. Figure 4As shown in b1, the upper half of the grain boundary consumes subgrains and grows due to growth advantage, therefore no obvious subgrain structure is observed above the grain. In contrast, the subgrain structure in the lower half of the grain boundary is adjacent to the grain boundary and does not have growth advantage. This evidence suggests that in the 7.5Co alloy, while there exists a highly efficient abnormal grain growth mechanism induced by subgrain anomalous growth, there also exists a less efficient abnormal grain growth mechanism induced by the subgrain low-density region. This is because the subgrain low-density region of the 7.5Co alloy is relatively narrow (…). Figure 3 (b) Grain boundaries can still cross the low-density subgrain region and undergo anomalous growth, resulting in a large number of anomalously grown grains and affecting the final grain size. In contrast,... Figure 4 In c, no subgrain structure was observed within the large anomalous grains, and the grain size was uniform with no obvious small grains detected; this indicates that almost all anomalous grains in the 7.5Co1Mo alloy are anomalous grains with high migration efficiency induced by subgrain anomalous growth. Combined with... Figure 3 The subgrain distribution characteristics in the 7.5Co1Mo alloy in c clearly show that the large width of the low-density subgrain region in the Co and Mo mixed alloy makes it difficult for the grain boundaries to receive high-energy subgrains from the grain center, thus suppressing the inefficient abnormal grain growth induced by the inhomogeneity of the low-density subgrain region. However, since there are still high-density, high-energy subgrains at the grain center, the abnormal subgrain growth process continues and results in high-efficiency abnormal grain growth in the alloy.

[0164] Test Example 6

[0165] The metallographic structures of the alloys obtained in Comparative Examples 11-15, Comparative Example 17, and Examples 2-5 were observed using a metallographic microscope. Based on the alloys of Comparative Examples 11 and 17, the changes in grain boundary migration distance of Fe-Mn-Al-Ni alloy (0Co alloy) and Fe-Mn-Al-Ni-Co-Mo alloy (7.5Co1Mo alloy) with increasing reheating time were recorded, and the curves are shown below. Figure 5 As shown in the figure, the starting point of abnormal grain growth is defined as the second heating temperature reaching 1180℃, and the average grain size at this point is taken as the starting point of grain boundary migration. The time in the figure is the sum of the heating time from 1180℃ and the holding time. As described in Test Example 5, since the two alloys exhibit different abnormal grain growth behaviors, the grain boundary migration distance of Comparative Examples 11-15 (0Co alloy) represents the maximum distance between the grain boundary of the residual subgrain region and the non-subgrain region within the abnormal grain. The grain boundary migration distance of Comparative Examples 17 and Examples 2-5 (7.5Co1Mo alloy) is the difference between the maximum grain size of the alloy and the average grain size at 1180℃.

[0166] from Figure 5It can be seen that the 7.5Co1Mo alloy exhibits the highest slope initially, but begins to decrease after reaching the maximum point. This is because initially, the number of abnormally grown grains is small, and the grain boundaries have sufficient space to migrate. However, after holding at temperature for a period of time, more grains undergo abnormal growth, and the abnormal grain growth is nearing completion. When two abnormal grains collide, the grain boundary migration on that side stops. Therefore, the maximum slope is taken as the growth rate of abnormal grains to eliminate this interference. Thus, the maximum grain boundary migration rate of the 7.5Co1Mo alloy is 21.1 × 10⁻⁶. -6 The grain boundary migration rate of the 0Co alloy is 2.8 × 10 m / s. -6 m / s, an increase of an order of magnitude. According to Figure 4 This is because the abnormal grain growth induced by the low-density subgrain region in the Co and Mo alloy is suppressed. Therefore, the abnormal grains in the 7.5Co1Mo alloy are all caused by the abnormal growth of subgrains. The Co and Mo elements synergistically improve the grain growth efficiency of the Fe-Mn-Al-Ni alloy, allowing the grains to grow faster.

[0167] Test Example 7

[0168] The alloys obtained in Comparative Example 16 and Example 6 were observed. In the alloy of Comparative Example 16, white grain boundaries were clearly visible. These grain boundaries intersected to form triangular grain boundaries, and more than ten grains were detected. In contrast, no white grain boundaries were detected in the alloy of Example 6, indicating that there was only one grain. Figure 4 Because the anomalous grain growth induced by the subgrain low-density region is suppressed in the Co and Mo alloys, this mechanism cannot generate anomalous grains, thus reducing the number of anomalous grains. As a result, the 7.5Co1Mo alloy (Example 6) reaches a single-crystal structure after four cooling-reheating cycles. Therefore, Co and Mo elements synergistically improve the grain growth capability of Fe-Mn-Al-Ni alloys, allowing the grains to grow larger.

[0169] Test Example 8

[0170] After removing the oxide scale from the hot-rolled billets of Comparative Example 15 and Example 5, 3mm × 3mm × 6mm specimens were cut, vacuum-sealed, and then subjected to cyclic heat treatment (without changing parameters) to obtain specimens. Comparative Example 15 specimens were subjected to cyclic loading and unloading compression tests using a tensile testing machine at a speed of 0.5mm / min. The loading procedure was as follows: first, load to 3%, then unload; then, gradually increase the strain by 2% and unload again until the specimen exhibited a large irreversible strain, and record the maximum loading strain at this point. Example 5 specimens were subjected to compression tests using a tensile testing machine at a loading speed of 0.5mm / min, directly loading to the maximum loading strain and then unloading. Specimen deformation was measured using a dial indicator with a preload stress of 50N. The recoverable strain for each load was calculated by subtracting the irreversible strain from the maximum loading strain for that load, resulting in the stress-strain curves shown below. Figure 6 As shown. From Figure 6 As can be seen, the 0Co alloy (Comparative Example 15) exhibits only ~3.5% of maximum recoverable strain, while the 7.5Co1Mo alloy (Example 5) exhibits a maximum recoverable strain as high as ~7.2%. This indicates that Co and Mo elements effectively improve the superelasticity of the alloy by enhancing grain growth efficiency and capability.

[0171] The above description is only a preferred embodiment of the present invention. It should be noted that for those skilled in the art, several improvements and modifications can be made without departing from the principle of the present invention, and these improvements and modifications should also be considered within the scope of protection of the present invention.

Claims

1. A Fe-Mn-Al-Ni based superelastic alloy, characterized in that, On an atomic percentage basis, it includes: Mn 34%, Al 15%, Ni 7.5%, Co 2~7.5%, Mo 0.3~1%, and the balance Fe; The preparation method of the Fe-Mn-Al-Ni based superelastic alloy includes the following steps: The Fe-Mn-Al-Ni-Co-Mo alloy castings were homogenized and then hot-rolled to obtain billets. The billet is subjected to cyclic heat treatment or directional recrystallization to obtain Fe-Mn-Al-Ni based superelastic alloy; The cyclic heat treatment includes the following steps: The billet is first heated to a first temperature, then cooled to a second temperature, and then heated to a third temperature. After the second temperature is raised to the third temperature, it is cooled to room temperature, or the cooling and second temperature-raising process is repeated and then cooled to room temperature. The first temperature is 1200~1300℃; the second temperature is not higher than 1180℃; the third temperature is 1190~1200℃; The parameters for directional recrystallization include: a hot zone temperature of 1180~1300℃, a hot zone width of 1~100mm, a temperature gradient of 10~500℃ / mm, and a pulling speed of 0.1~1000μm / s.

2. The Fe-Mn-Al-Ni based superelastic alloy according to claim 1, characterized in that, On an atomic percentage basis, it includes: Mn 34%, Al 15%, Ni 7.5%, Co 5~7.5%, Mo 0.4~1%, and the balance Fe.

3. The method for preparing a Fe-Mn-Al-Ni based superelastic alloy according to claim 1 or 2, characterized in that, Includes the following steps: The Fe-Mn-Al-Ni-Co-Mo alloy castings were homogenized and then hot-rolled to obtain billets. The billet is subjected to cyclic heat treatment or directional recrystallization to obtain Fe-Mn-Al-Ni based superelastic alloy.

4. The preparation method according to claim 3, characterized in that, The homogenization treatment is carried out at a temperature of 1200~1300℃ for 1~10h.

5. The preparation method according to claim 3, characterized in that, The initial rolling temperature of the hot rolling is not lower than 1200℃, and the final rolling temperature of the hot rolling is not lower than 1100℃.

6. The preparation method according to claim 3 or 5, characterized in that, The deformation amount of the hot rolling is not less than 50%.

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