A method for manufacturing a phase change alloy single crystal without deformation treatment, a phase change alloy single crystal, a phase change alloy single crystal member, and a terminal product

Abstract

This application relates to the field of metal material preparation technology, and discloses a method for preparing phase change alloy single crystals without deformation treatment. The method includes: depositing a first alloy layer by layer using electron beam filament deposition and performing interlayer remelting to obtain a first intermediate; wherein the first alloy exhibits reversible martensitic phase transformation characteristics; subjecting the first intermediate to thermal cycling treatment to obtain a second intermediate; and subjecting the second intermediate to temperature gradient annealing to obtain the first alloy single crystal. This preparation method, through the synergistic process of electron beam filament deposition, interlayer remelting, thermal cycling treatment, and temperature gradient annealing, eliminates the need for pre-deformation treatment, enabling efficient direct preparation of bulk phase change alloy single crystals from the solidified state. This achieves near-net-shape rapid prototyping of complex single crystal components and is suitable for the rapid preparation of large-size, complex single crystal components of phase change alloys. This application also discloses a phase change alloy single crystal, a phase change alloy single crystal component, and an end product.

Description

Technical Field

[0001] This application relates to the field of metal material preparation technology, such as a method for preparing phase change alloy single crystals without deformation treatment, phase change alloy single crystals, phase change alloy single crystal components and end products. Background Technology

[0002] NiTi alloys exhibit excellent shape memory, superelasticity, high damping, and elastothermal effects due to temperature / stress-induced reversible thermoelastic martensitic phase transformation. They also possess good wear and corrosion resistance and biocompatibility, making them widely used in aerospace, biomedicine, and solid-state refrigeration. The properties of NiTi alloys are strongly dependent on grain size; however, plastic deformation often occurs at grain boundaries in polycrystalline structures, leading to stress concentration and hindering reversible martensitic phase transformation, resulting in functional degradation. In contrast, NiTi single crystals offer superior phase transformation controllability and functional stability, simultaneously exhibiting large fully recoverable strain, high energy density, and ultra-long fatigue life. Furthermore, single-oriented single crystals eliminate grain boundary interference on the phase transformation path, providing accurate phase transformation thermo / kinetic data, which is crucial for a deeper understanding of material phase transformation behavior and the establishment of cross-scale constitutive models.

[0003] Existing single crystal preparation technologies have significant limitations. Traditional melt methods (such as Czochralski, zone melting, and directional solidification) require the material to be in a molten state, which easily introduces crucible contamination, involves complex and expensive equipment, consumes a lot of energy, and is inefficient, making it difficult to prepare large-size bulk single crystals. Existing solid-state single crystal growth methods often require pre-deformation treatment. For example, strain annealing requires cold rolling (deformation >70%) to introduce dislocation substructures, cyclic heat treatment also requires hot rolling / cold rolling to introduce high-density dislocations, and texture engineering grain growth methods rely on cold drawing deformation to introduce strong texture, providing driving force and structural basis for subsequent abnormal grain growth.

[0004] However, the pre-deformation process is complex, grain orientation control is difficult, and it is only applicable to simple-shaped single-crystal samples. Furthermore, NiTi alloys exhibit poor melting zone stability, high recrystallization activation energy, and low phase transformation temperature (<100℃), making it difficult to achieve bulk single-crystal growth with a single orientation using the aforementioned preparation techniques. Simultaneously, this alloy has a high work hardening index and severe springback, making it difficult to machine into complex structural parts. Therefore, there is an urgent need to develop a preparation technique that does not require pre-deformation and can directly obtain bulk NiTi single crystals from the solidified state.

[0005] It should be noted that the information disclosed in the background section above is only used to enhance the understanding of the background of this application, and therefore may include information that does not constitute prior art known to those skilled in the art. Summary of the Invention

[0006] To provide a basic understanding of some aspects of the disclosed embodiments, a brief summary is given below. This summary is not intended as a general commentary, nor is it intended to identify key / important components or describe the scope of protection of these embodiments, but rather as a prelude to the detailed description that follows.

[0007] This disclosure provides a method for preparing phase change alloy single crystals without deformation treatment, phase change alloy single crystals, phase change alloy single crystal components, and end products, in order to solve the technical problem that existing processes are difficult to achieve the preparation of bulk single crystals of phase change alloys with a single orientation, and that the grain size of the single crystals is small.

[0008] In some embodiments, the method for preparing a single crystal of a phase transformation alloy without deformation treatment is characterized by comprising: depositing a first alloy layer by layer using an electron beam fused wire deposition process and performing interlayer remelting to obtain a first intermediate; wherein the first alloy has reversible martensitic phase transformation characteristics; subjecting the first intermediate to thermal cycling treatment to obtain a second intermediate; and subjecting the second intermediate to temperature gradient annealing to obtain a single crystal of the first alloy.

[0009] In some embodiments, the phase change alloy single crystal is prepared using the method for preparing a phase change alloy single crystal without deformation treatment as described in any of the foregoing embodiments.

[0010] In some embodiments, the phase change alloy single crystal component is composed of a first alloy single crystal, which is prepared by the method for preparing a phase change alloy single crystal without deformation treatment as described in any of the foregoing embodiments.

[0011] In some embodiments, the end product includes a phase change alloy single crystal as described above, or includes a phase change alloy single crystal component as described above.

[0012] The present disclosure provides a method for preparing a phase change alloy single crystal without deformation treatment, a phase change alloy single crystal, a phase change alloy single crystal component, and a terminal product, which can achieve the following technical effects: By employing a repetitive layer-by-layer deposition-interlayer remelting strategy, crystal growth orientation can be effectively controlled. High-density dislocations are introduced through thermal cycling, and combined with temperature gradient annealing, directional growth of single crystals can be achieved, ultimately forming millimeter-scale bulk single crystal materials with high orientation consistency. This preparation method, through the synergistic process of electron beam filament deposition, interlayer remelting, thermal cycling, and temperature gradient annealing, eliminates the need for pre-deformation treatment and can directly and efficiently prepare bulk single crystals of phase change alloys from the solidified state. This enables near-net-shape rapid prototyping of complex single crystal components and is suitable for the rapid fabrication of large-size, complex single crystal components of phase change alloys.

[0013] The orientation of single-crystal phase transformation alloys is precisely controllable, with no grain boundary interference within the crystal, resulting in a uniform and stable phase transformation process. The shape memory effect recovery rate and superelastic properties are significantly better than those of polycrystalline and deformed single-crystal materials, and the phase transformation temperature fluctuation range is narrow.

[0014] The phase change alloy single crystal component is composed of the first alloy single crystal as the core and is prepared by a synergistic process of electron beam filament deposition pretreatment, thermal cycling treatment and temperature gradient annealing. It has high density, high orientation consistency and low defect density, good shape memory effect and superelasticity, and strong structural stability.

[0015] The above general description and the description below are exemplary and illustrative only and are not intended to limit this application. Attached Figure Description

[0016] One or more embodiments are illustrated by way of example with reference to the accompanying drawings. These illustrations and drawings do not constitute a limitation on the embodiments. Elements having the same reference numerals in the drawings are shown as similar elements. The drawings are not to be scaled. And wherein: Figure 1 This is a flowchart illustrating a method for preparing a single crystal of a phase change alloy without deformation, as provided in an embodiment of this disclosure. Figure 2 This is a flowchart of another method for preparing a phase change alloy single crystal without deformation treatment provided in this embodiment of the disclosure; Figure 3 This is a flowchart of another method for preparing a phase change alloy single crystal without deformation treatment provided in this embodiment of the disclosure; Figure 4 This is a flowchart of another method for preparing a phase change alloy single crystal without deformation treatment provided in this embodiment of the disclosure; Figure 5 This is a flowchart of another method for preparing a phase change alloy single crystal without deformation treatment provided in this embodiment of the disclosure; Figure 6 This is a schematic diagram of an electron beam filament deposition coupled interlayer remelting technique for a method of preparing a single crystal of a phase change alloy without deformation treatment provided in this embodiment of the present disclosure; Figure 7 This is a schematic diagram of the thermal cycling process of a method for preparing a single crystal of a phase change alloy without deformation provided in this embodiment of the present disclosure; Figure 8 This is a temperature gradient annealing schematic diagram of a method for preparing a single crystal of a phase change alloy without deformation provided in this embodiment of the present disclosure; Figure 9 This is a crystallographic texture distribution diagram of an internal sample of a first intermediate Ti-50.7at.%Ni alloy provided in this embodiment of the present disclosure; Figure 10 This is the XRD pattern of a Ti-50.7at.%Ni alloy sample, a first intermediate, provided in an embodiment of this disclosure; Figure 11 This is a schematic diagram of the evolution of the internal microstructure of a Ti-50.7at.%Ni alloy sample, which is the second intermediate in a method for preparing a single crystal of a phase transformation alloy without deformation treatment provided in this embodiment of the present disclosure. Figure 12 This is an internal microstructure orientation diagram of a Ti-50.7at.%Ni alloy single crystal, provided by an embodiment of this disclosure for a method of preparing a phase transformation alloy single crystal without deformation treatment. The diagram on the left is the corresponding {001} pole figure, and the diagram on the right is the standard inverse pole figure of B2 austenite. Figure 13 The phase spectrum of the gradient NiTi alloy prepared in Comparative Example 1 is shown. Figure 14 This is the solidification orientation diagram of the gradient NiTi alloy prepared in Comparative Example 1.

[0017] Figure label: 10: Substrate; 20: Deposited layer; 201: Molten pool; 30: Electron beam; 50: First alloy; 60: First intermediate; 601: B2 austenite; 602: B19′ martensite; 70: Second intermediate; 701: Single crystal region; 702: Columnar crystal region; 801: Thermal environment; 802: Cold environment; 90: Dual-temperature zone device; 901: Furnace chamber; 902: High-temperature zone; 903: Low-temperature zone. Detailed Implementation

[0018] To provide a more detailed understanding of the features and technical content of the embodiments of this disclosure, the implementation of the embodiments of this disclosure will be described in detail below with reference to the accompanying drawings. The accompanying drawings are for illustrative purposes only and are not intended to limit the embodiments of this disclosure. In the following technical description, for ease of explanation, several details are used to provide a full understanding of the disclosed embodiments. However, one or more embodiments may still be implemented without these details. In other cases, well-known structures, steps, and apparatus may be simplified to simplify the illustrations.

[0019] The terms "first," "second," etc., used in the specification, claims, and accompanying drawings of this disclosure are used to distinguish similar objects and are not necessarily used to describe a specific order or sequence. It should be understood that such terms can be used interchangeably where appropriate for the embodiments of this disclosure described herein. Furthermore, the terms "comprising" and "having," and any variations thereof, are intended to cover non-exclusive inclusion.

[0020] Unless otherwise defined, all technical and scientific terms used herein have the same meaning as commonly understood by one of ordinary skill in the art to which this application belongs. The term "multiple" means two or more. In embodiments of this disclosure, the character " / " indicates an "or" relationship between the preceding and following objects. For example, A / B means: A or B. The term "and / or" describes an association between objects, indicating that three relationships can exist. For example, A and / or B means: A or B, or, A and B.

[0021] It should be noted that, unless otherwise specified, the embodiments and features described in the present disclosure can be combined with each other.

[0022] Those skilled in the art will understand that in the methods described in this application and other parts thereof, for example, in the methods of various embodiments, examples, or claims, the order in which the steps are written does not imply a strict execution order and does not constitute any limitation on the implementation process. The detailed execution order of each step should be determined by its function and possible internal logic. Unless otherwise specified, all steps of this application may be performed sequentially or randomly, preferably sequentially. For example, if the method includes steps A and B, it means that the method may include steps A and B performed sequentially, or it may include steps B and A performed sequentially. For example, if the method may also include step C, it means that step C can be added to the method in any order. For example, the method may include steps A, B, and C, or it may include steps A, C, and B, or it may include steps C, A, and B, etc.

[0023] In this application, open-ended technical features or solutions described using terms such as "containing," "including," or "comprising" do not exclude additional members beyond those listed unless otherwise specified. They can be considered as providing both closed-ended features or solutions comprised of the listed members and open-ended features or solutions that include additional members beyond the listed members. For example, if A includes a1, a2, and a3, it may also include other members or exclude additional members unless otherwise specified. This can be considered as providing both the feature or solution that "A consists of a1, a2, and a3" and the feature or solution that "A includes not only a1, a2, and a3, but also other members."

[0024] The "range" disclosed in this application can be defined in the form of a lower limit and an upper limit. A given range is defined by selecting a lower limit and an upper limit, which define the boundaries of the specific range. Ranges defined in this way can include or exclude endpoints; any endpoint can be independently included or excluded, and they can be arbitrarily combined, meaning any lower limit can be combined with any upper limit to form a range. For example, if ranges of 30~100 and 80~90 are listed for a specific parameter, it is expected that ranges of 30~90 and 80~100 will also be included. Furthermore, if minimum range values ​​of 1 and 2 are listed, and if maximum range values ​​of 3, 4, and 5 are also listed, then the following ranges can all be expected: 1~3, 1~4, 1~5, 2~3, 2~4, and 2~5. In this application, unless otherwise stated, the numerical range "a~b" represents a shortened representation of any combination of real numbers between a and b, where a and b are real numbers. For example, the numerical range "0~7" indicates that all real numbers between "0" and "7" have been listed in this article; "0~7" is simply a shortened representation of these numerical combinations. Furthermore, when a parameter is described as an integer ≥2, it is equivalent to listing integers such as 2, 3, 4, 5, 6, 7, 8, 9, 10, 11, 12, etc. For instance, when a parameter is described as an integer selected from "2~10", it is equivalent to listing the integers 2, 3, 4, 5, 6, 7, 8, 9, and 10.

[0025] In this application, unless otherwise specified, A (e.g., B) means that B is a non-limiting example of A, and it is understood that A is not limited to B.

[0026] This disclosure provides a method for preparing phase transformation alloy single crystals without deformation treatment, including: S110. Using electron beam fused wire deposition process, the first alloy is deposited layer by layer and remelted between layers to obtain the first intermediate; wherein, the first alloy has reversible martensitic phase transformation characteristics.

[0027] S120. The first intermediate is subjected to thermal cycling treatment to obtain the second intermediate.

[0028] S130. The second intermediate is subjected to temperature gradient annealing to obtain the first alloy single crystal.

[0029] Electron beam fused deposition coupled with interlayer remelting technology enables precise melting and layer-by-layer epitaxial formation of phase transformation alloys. This effectively refines the initial grains, eliminates interlayer porosity and incomplete fusion defects, and constructs an initial microstructure with a <001> preferred orientation along the deposition direction, providing a high-quality seed crystal foundation for subsequent single crystal growth. Simultaneously, the electron beam fused deposition process provides a high-vacuum environment, effectively isolating impurities such as oxygen and nitrogen, reducing the formation of brittle phases such as oxide and nitride phases, and effectively improving the purity and compositional uniformity of the first intermediate 60. Furthermore, leveraging the additive manufacturing characteristics of electron beam fused deposition, near-net-shape forming of large-sized components can be directly achieved, providing dimensional support for bulk single crystal fabrication. Bulk single crystals refer to single crystals with grain sizes reaching the millimeter level.

[0030] For the reversible martensitic phase transformation characteristics of the first alloy 50, thermal cycling treatment can trigger a complete reversible phase transformation from austenite to martensite to austenite. Introducing high-density dislocations through phase transformation induction can replace the dislocation introduction function of traditional plastic deformations such as cold rolling and hot rolling, providing sufficient energy driving force for grain growth. The entire process requires no deformation operations, which not only simplifies the process flow but also fundamentally avoids defects such as processing cracks and stress concentration caused by deformation, thereby effectively improving the structural integrity of the first intermediate 60.

[0031] Temperature gradient annealing guides grain boundary migration through a directional temperature field, driving the directional growth of <001> dominant orientation grains and their engulfment of other orientation grains, ultimately forming a grain boundary-free, highly oriented bulk single crystal of phase transformation alloy.

[0032] The method for preparing phase change alloy single crystals without deformation treatment provided in this disclosure effectively controls crystal growth orientation through a repeated layer-by-layer deposition-interlayer remelting strategy. High-density dislocations are introduced using thermal cycling, and combined with temperature gradient annealing, directional growth of the single crystal is achieved, ultimately forming a millimeter-scale bulk single crystal material with high orientation consistency. This preparation method, through the synergistic process of electron beam filament deposition, interlayer remelting, thermal cycling, and temperature gradient annealing, eliminates the need for pre-deformation treatment and can directly and efficiently prepare bulk phase change alloy single crystals from the solidified state. This enables near-net-shape rapid prototyping of complex single crystal components and is suitable for the rapid fabrication of large-size, complex single crystal components of phase change alloys.

[0033] Optionally, the first alloy 50 uses wire as the deposition material.

[0034] The first alloy 50 uses wire as the deposition material, which is compatible with the electron beam fused wire deposition process, making the feeding continuous and stable, and improving the consistency of forming.

[0035] The core parameters of electron beam fused wire deposition (EBD) include accelerating voltage, focusing current, wire feed speed, printing speed, first beam current density, and vacuum level. These process parameters can be flexibly adjusted according to actual preparation needs. By matching differentiated parameter combinations for deposition and interlayer remelting, repeated layer-by-layer deposition and interlayer remelting of the first alloy can be achieved, ensuring forming stability and microstructure quality.

[0036] Optionally, the process parameters used in the electron beam fused deposition process include: an accelerating voltage of 55~65 kV. Optionally, the accelerating voltage is 60 kV.

[0037] Optionally, the focusing current is 950~1050 mA. Optionally, the focusing current is 1000 mA.

[0038] Optionally, the wire feeding speed is 2500~3500 mm / min. Optionally, the wire feeding speed is 3000 mm / min.

[0039] Optionally, the printing speed is 550~650 mm / min. Optionally, the printing speed is 600 mm / min.

[0040] Optionally, the first beam current density is 45~50 mA. Optionally, the first beam current density is 45 mA.

[0041] Optionally, the vacuum level is less than or equal to 7 × 10⁻⁶. -2 Pa. Optionally, the vacuum degree is 7 × 10⁻⁶. -2 Pa.

[0042] The above combination of process parameters is precisely matched with the electron beam fused wire deposition process, ensuring complete melting of the first alloy wire and layer-by-layer epitaxial formation. The synergistic effect of the accelerating voltage of 55–65 kV, focusing current of 950–1050 mA, and first beam current density of 45–50 mA allows for precise control of the molten pool energy and size. The wire feed speed of 2500–3500 mm / min is matched with the printing speed of 550–650 mm / min, guaranteeing deposition efficiency and forming stability. The vacuum level of the high-vacuum environment is less than or equal to 7 × 10⁻⁶. - ²Pa can reduce the intrusion of oxygen and nitrogen impurities and inhibit the formation of brittle phases. This combination of parameters can effectively improve the density and compositional uniformity of the first alloy single crystal, and give the first alloy single crystal an initial microstructure with a strong <001> preferred orientation.

[0043] In some embodiments, in step S110, an electron beam fused wire deposition process is used to deposit the first alloy layer by layer and perform interlayer remelting to obtain a first intermediate 60, comprising: S111. An electron beam is used to deposit a first alloy on the substrate in a unidirectional scanning mode at a first beam current density to form a phase change alloy deposition layer.

[0044] S112. Using an electron beam at a second beam current density, the current phase change alloy deposited layer is remelted to form a molten pool on the surface of the current phase change alloy deposited layer.

[0045] S113, A first alloy layer is deposited on the surface of the molten pool.

[0046] S114. Perform a deposition-remelting cycle to obtain the first intermediate.

[0047] In the electron beam fused deposition process, an electron beam 301 bombards the first alloy 50 and the substrate 10 with high energy density, causing the first alloy to melt instantly and partially fuse with the surface layer of the substrate 10, forming a continuous and stable liquid metal pool 201. This pool 201 remains stable in a high vacuum environment, serving as both the core area for melting the first alloy and a key carrier for achieving layer-by-layer deposition and solidification.

[0048] Electron beam fused deposition is employed, with the first alloy 50 deposited layer by layer onto the substrate 10 in a unidirectional scanning mode. After each layer is deposited, the current layer is remelted at a second beam current density. This deposition-remelting cycle is repeated to obtain the first intermediate 60. The deposition-remelting cycle is performed sequentially as follows: First, a phase change alloy deposition layer 20 is deposited. After the layer has initially solidified, an electron beam 301 is used to precisely bombard the surface of the current phase change alloy deposition layer 20. By controlling the beam current density and scanning speed of the electron beam 301, the surface of the deposition layer 20 undergoes controllable melting, forming a uniform molten pool 201. Immediately afterwards, a next phase change alloy layer is deposited on the surface of the molten pool 201. At this point, the molten pool 201 serves as an interface transition zone between the two deposition layers, promoting atomic diffusion and metallurgical bonding between adjacent layers. The above deposition-remelting operation is repeated until the first intermediate 60 of a preset size is obtained, effectively eliminating defects such as interlayer porosity and cracks.

[0049] The unidirectional scanning mode enables the molten pool 201 to form a consistent heat dissipation gradient along the scanning direction during solidification, inducing grains to grow in a single direction and initially forming a texture structure with <001> as the preferred orientation.

[0050] Unidirectional scanning mode refers to scanning in the same direction for each layer. Unidirectional scanning constructs a stable directional heat dissipation gradient, driving grains to grow along the optimal heat conduction direction. Combined with interlayer remelting to optimize texture consistency, a strong <001> preferred orientation is formed, providing a high-quality seed crystal foundation for subsequent single crystal growth. Interlayer remelting effectively refines grains, fills interlayer porosity and incomplete fusion defects, and, combined with the high vacuum environment of electron beam fused deposition, reduces the intrusion of oxygen and nitrogen impurities and the formation of brittle phases, effectively improving the density of the first intermediate 60. Repeated deposition-remelting cycles ensure the stability of the forming quality of each layer and leverage additive manufacturing characteristics to achieve near-net-shape forming of large-sized, complex-shaped intermediates. Simultaneously, it avoids deformation defects inherent in traditional pretreatment, simplifying the process flow. Through the synergy of unidirectional scanning deposition and interlayer remelting cycles, the first intermediate 60 possesses the core advantages of uniform texture, dense quality, high purity, and large size, providing a microstructure basis for subsequent thermal cycling and temperature gradient annealing.

[0051] Optionally, the second beam current density is 20~25 mA. Optionally, the second beam current density is 20 mA.

[0052] A second beam current density of 20–25 mA falls within the low-energy beam range, effectively remelting the deposited layer to refine grains, fill interlayer porosity and incomplete fusion defects, without over-melting and damaging the formed structure. When the second beam current density is greater than or equal to 20 mA, insufficient remelting energy is avoided, effectively eliminating micropores and grain boundary segregation within the deposited layer. When the second beam current density is less than or equal to 25 mA, excess energy is avoided, preventing excessive depth of the molten pool 201 from inducing interlayer mixing and protecting the initial texture orientation. A second beam current density of 20 mA achieves a balance between defect elimination and texture preservation, more effectively controlling the depth of the molten pool 201 and suppressing the formation of brittle phases.

[0053] Optionally, in step S113, the deposition-remelting cycle is repeated until a first height is reached to obtain a first intermediate 60, the first height being 20~40 mm. Optionally, the first height is 30 mm.

[0054] A first height of 20–40 mm balances single-crystal growth efficiency and quality, aligning with the large-size forming advantages of electron beam fused deposition modeling. When the first height is greater than or equal to 20 mm, insufficient space for grain boundary migration during subsequent annealing is avoided, facilitating the formation of bulk single crystals. When the first height is less than or equal to 40 mm, a uniform temperature gradient within the first intermediate is achieved, preventing initial grain orientation dispersion. A first height of 30 mm avoids insufficient space for subsequent process control due to excessively low height, while also preventing heat accumulation and stress concentration caused by excessively high height, thus meeting the grain boundary migration requirements of subsequent temperature gradient annealing.

[0055] The synergy of repeated deposition-remelting cycles and parameter combinations ensures uniform microstructure quality in each layer, and large-size intermediates meeting requirements can be obtained without additional pretreatment. Low-energy beam remelting and reasonable height design reduce the risk of deformation and cracking during the forming process. Combined with the texture control effect of unidirectional scanning, a strong <001> preferred orientation is formed.

[0056] Optionally, the substrate 10 is a metal substrate.

[0057] Metal substrates possess excellent thermal conductivity and structural rigidity, enabling rapid heat conduction from the molten pool during electron beam filament deposition. This prevents uncontrolled molten pool 201 or abnormal grain growth caused by excessively high local temperatures, ensuring dimensional accuracy and forming stability of layer-by-layer deposition.

[0058] Optionally, the substrate is a TA1 pure titanium substrate.

[0059] TA1 pure titanium substrates exhibit excellent thermal conductivity matching, with a coefficient of thermal expansion that is relatively close to that of alloys such as NiTi, thus preventing interfacial cracking caused by thermal stress during deposition. Furthermore, TA11 titanium alloy substrates possess excellent corrosion resistance, making them suitable for demanding operating conditions. It is understandable that the substrate can also be a TA11 pure titanium substrate.

[0060] In some embodiments, the second beam density of the interlayer remelting is lower than the first beam density of the deposition stage.

[0061] Interlayer remelting employs a second beam density lower than that of the deposition stage. This process can refine grains, fill interlayer pores and unfused defects through moderate remelting, reduce the intrusion of oxygen and nitrogen impurities and the formation of brittle phases, and improve the density and compositional uniformity of the first intermediate. It can also avoid over-melting and destroying the already formed directional structure, continue the stable directional heat dissipation gradient constructed by unidirectional scanning, strengthen the coherence of the <001> preferred orientation, and provide a high-quality seed crystal foundation for subsequent single crystal growth. At the same time, in conjunction with repeated deposition-remelting cycles, it can achieve stable near-net-shape forming of large-size and complex-shaped intermediates by relying on additive manufacturing characteristics.

[0062] Optionally, the remelting scan path is consistent with the deposition path. This avoids disrupting the orientation consistency of the already deposited layers.

[0063] In some embodiments, in step S120, the first intermediate is subjected to thermal cycling to obtain the second intermediate 70, including: S121. The first intermediate is alternately immersed in a hot environment at a first temperature and a cold environment at a second temperature to form a two-way cycle of hot to cold and cold to hot; wherein the first temperature is higher than the austenitic phase transformation end temperature of the first alloy, and the second temperature is lower than the martensitic phase transformation end temperature of the first alloy.

[0064] The first intermediate 60 is immersed in a hot environment 801 at a first temperature, a cold environment 802 at a second temperature, and then in a hot environment 801 at the first temperature in sequence. By immersing in these environments alternately, a two-way cycle from hot to cold and from cold to hot is formed.

[0065] When the first temperature is higher than the austenitic transformation end temperature (Af), the first alloy can completely transform into the austenitic phase, with the atomic arrangement becoming regularized, eliminating the lattice distortion generated during the deposition process. When the second temperature is lower than the martensitic transformation end temperature (Mf), the first alloy can completely transform into the martensitic phase, with atoms sliding along specific crystal planes during the transformation, forming a large number of dislocations.

[0066] This complete austenitization-martenitization cycle can generate a high-density dislocation substructure without plastic deformation. High-density dislocations reduce the activation energy for grain boundary migration. During subsequent temperature gradient annealing, grain boundaries migrate along directions with lower dislocation density. Uniformly distributed dislocations guide grain boundaries to move in a single direction, avoiding the polycrystalline mixing problem caused by random grain boundary migration in the absence of dislocations. Furthermore, the alternating soaking method, compared to continuous heating or cooling, allows for more precise control of the phase transformation, avoiding uneven dislocation distribution caused by incomplete local phase transformation. This effectively improves the uniformity of the second intermediate 70 microstructure, providing a directional grain boundary driving force for the subsequent preparation of single-oriented single crystals.

[0067] This thermal cycling process, through alternating immersion at a first temperature above the austenitic transformation end temperature and a second temperature below the martensitic transformation end temperature, drives the first intermediate 60 to repeatedly undergo reversible austenitic-martensite transformations. For example, when the transformation alloy is a NiTi alloy, the formed first intermediate 60 can repeatedly undergo reversible transformations from B2 austenite 601 to B19' martensite. This induces the formation of uniformly distributed defects such as dislocations, providing active sites for directional grain growth in subsequent temperature gradient annealing, and further eliminates residual component segregation, releases residual stress generated during deposition, and reduces the formation of brittle phases. Simultaneously, it ensures precise and stable atomic ratios, enabling the second intermediate 70 to possess both high density, uniform microstructure, and strong <001> preferred orientation.

[0068] Optionally, the thermal cycling process is a bidirectional cycle of thermal environment 801, cold environment 802, and thermal environment 801, and in each cycle, the immersion time of thermal environment 801 is the same as the immersion time of cold environment 802.

[0069] Immersion in hot environment 801 and cold environment 802 at equal times ensures that the first intermediate 60 completes the full austenitization-martenitization-reaustitization phase transformation cycle in each cycle. This symmetrical thermal cycling mode avoids incomplete local phase transformation caused by uneven immersion time. When the first alloy is a NiTi alloy, this allows for repeated B2 formation within the NiTi alloy. The B19′ reversible phase transformation uniformly induces the generation of defects such as dislocations, and the defect distribution has no obvious clustering areas, providing globally uniform active sites for the directional growth of grains in subsequent temperature gradient annealing, and enhancing the coherence of the <001> preferred orientation.

[0070] In some embodiments, the thermal environment 801 is a water bath environment. Optionally, the first temperature is 40~100 ℃. Optionally, the first temperature is 90 ℃. Optionally, the immersion time of the thermal environment 801 each time is 2~5 min. Optionally, the immersion time of the thermal environment 801 each time is 3 min.

[0071] Optionally, the cold environment 802 is a liquid nitrogen environment. Optionally, the second temperature is -196~-30 ℃. Optionally, the second temperature is -196 ℃. Optionally, the immersion time in the cold environment 802 each time is 2~5 min. Optionally, the immersion time in the cold environment 802 each time is 3 min.

[0072] Optionally, the alternating soaking cycle is repeated 40 to 80 times. Optionally, the alternating soaking cycle is repeated 50 times.

[0073] The first temperature of the thermal environment 801 is set to 40~100 ℃, higher than the austenitic transformation end temperature of the transformation alloy, ensuring complete austenitization of the first intermediate. The second temperature of the cold environment 802 is set to -196~-30 ℃, lower than the martensitic transformation end temperature of the transformation alloy, ensuring complete martensitization of the first intermediate and ultimately achieving a complete reversible phase transformation of the first intermediate. For example, the austenitic transformation end temperature of NiTi alloy is 25~35 ℃, and the martensitic transformation end temperature of NiTi alloy is -25~-15 ℃, thus enabling the formation of the B2 intermediate in NiTi alloy. B19′ complete reversible phase transition.

[0074] The first intermediate is immersed in the hot environment 801 and the cold environment 802 for 2-5 minutes, which can avoid incomplete local phase transformation caused by uneven immersion time. Combined with 30-80 cycles, the alloy interior repeatedly undergoes a complete phase transformation cycle. This can uniformly induce the generation of dislocation defects without obvious aggregation regions, providing globally uniform active sites for directional grain growth.

[0075] Meanwhile, the intense thermal expansion and contraction effect generated by the extreme temperature difference between the thermal environment 801 and the cold environment 802 can quickly release the residual stress accumulated during electron beam fused deposition, making it particularly suitable for large-sized first intermediates. A reasonable cycle count and isochronous design can balance defect induction and stress release, avoid secondary stress generation, and ensure the structural integrity and microstructure uniformity of the first intermediate.

[0076] The thermal cycling treatment uses a hot environment of 30~100 ℃ and a cold environment of -196~-20 ℃, and is immersed in hot and cold environments for 2~5 min for 30~80 times. This can match the phase transformation characteristics of the phase transformation alloy and efficiently control the microstructure defects and stress state of the first intermediate.

[0077] In some embodiments, in step S130, the second intermediate 70 is subjected to temperature gradient annealing to obtain a first alloy single crystal, including: S131. The second intermediate is placed in a dual-temperature zone device, the target temperature of the high-temperature zone is set as the third temperature, the target temperature of the low-temperature zone is set as the fourth temperature, the high-temperature zone and the low-temperature zone are heated to the target temperature at the first rate, and then kept at the temperature for a first time and cooled to room temperature.

[0078] Optionally, the third temperature is 1120~1200 ℃. Optionally, the third temperature is 1160 ℃.

[0079] Optionally, the fourth temperature is 1060~1140 ℃. Optionally, the fourth temperature is 1100 ℃.

[0080] Optionally, the first rate is 5~15 ℃ / min. Optionally, the first rate is 10 ℃ / min.

[0081] Optionally, the first duration is 48~96 hours. Optionally, the first duration is 72 hours.

[0082] Combination Figure 8 As shown, one end of the furnace 901 of the dual-temperature zone device 90 is designated as a high-temperature zone 902, and the other end as a low-temperature zone 903. A second intermediate 70 is placed inside the furnace 901 of the dual-temperature zone device 90, with one end of the second intermediate 70 positioned in the high-temperature zone 902 and the other end in the low-temperature zone 903. The end of the second intermediate 70 in the high-temperature zone 902 forms a single-crystal region 701, and the end in the low-temperature zone 903 forms a columnar crystal region 702. The temperature gradient direction from the high-temperature zone to the low-temperature zone in the furnace 901 is as follows: Figure 8As shown, a dual-temperature gradient annealing method is employed. A stable temperature difference between the high-temperature and low-temperature zones creates a static temperature gradient along the length of the second intermediate. Unlike the gradientless environment of traditional homogenous annealing, this static temperature gradient precisely guides grain boundaries to migrate directionally from the high-temperature zone to the low-temperature zone. This forces dominant oriented grains (such as the <001> orientation in NiTi alloys) to preferentially engulf other oriented grains, effectively avoiding the polycrystalline formation problem caused by random grain boundary migration during homogenous annealing and significantly improving the single-crystallization success rate.

[0083] Meanwhile, the directional heat flow precisely matches the strong <001> fiber texture of the second intermediate 70 along the construction direction, providing a continuous growth driving force for the dominant grains. Combined with the slow heating rate of the first rate and the long holding time of the first duration, it can ensure that the grain boundaries have sufficient migration activity, and suppress the nucleation of new grains and the abnormal growth of brittle phases, ultimately obtaining a high-quality alloy single crystal with high orientation consistency, low defect density and high density.

[0084] The third temperature in the high-temperature zone, ranging from 1120 to 1200 °C, matches the thermodynamic properties of the first alloy. When the third temperature is greater than or equal to 1120 °C, the first alloy can achieve complete recrystallization, imparting strong grain boundary migration capabilities; when the third temperature is less than or equal to 1200 °C, melting of the first alloy or abnormal grain growth can be avoided, ensuring microstructure uniformity. Taking NiTi alloy as an example, 1160 °C is the optimal recrystallization temperature for NiTi alloy, effectively increasing the grain boundary migration rate and accelerating the dominance of dominant grains.

[0085] The fourth temperature in the low-temperature region is set at 1060~1140 ℃, creating a temperature difference of 40~80 ℃ between the low-temperature and high-temperature regions, thus stabilizing the temperature gradient. This allows for a continuous and gentle heat flow direction, which aligns with the <001> fiber texture formed by electron beam fused deposition, providing directional growth momentum for predominantly oriented grains while suppressing new grain nucleation and reducing grain boundary interference. The stable temperature gradient design is suitable for large-size second intermediate 70, ensuring uniform single crystallization throughout the entire component and avoiding localized incomplete single crystallization.

[0086] A heating rate of 5–15 °C / min is adopted. When the first rate is greater than or equal to 5 °C, the process efficiency can be improved. When the first rate is less than or equal to 15 °C, the temperature change rate can be slowed down, avoiding thermal stress cracks in the second intermediate 70 due to thermal shock, while providing sufficient time for atomic diffusion and reducing component segregation. When the first heating rate is 10 °C / min, uniform heating can be achieved, which not only ensures the thermal environment 801 required for recrystallization and grain boundary migration, but also completely avoids thermal stress cracks, thus balancing process efficiency and component integrity.

[0087] The holding time is 48~96 h. When the first holding time is greater than or equal to 48 h, the grain boundaries can migrate fully, allowing the dominant grains to completely engulf other oriented grains and form a single grain. When the first holding time is less than or equal to 96 h, the energy waste caused by excessive holding can be avoided, and the process cost can be controlled. When the first holding time is 72 h, the dominant oriented grains (such as <001> orientation) can grow fully, eliminating residual grain boundaries and micro-defects, forming a complete single crystal that runs through the entire component, and suppressing the abnormal precipitation of brittle phases, thus ensuring the metallurgical purity of the single crystal.

[0088] The furnace cooling method can slowly release the residual stress accumulated during annealing, avoid phase transformation cracks caused by rapid cooling, maintain the orientation consistency of the single crystal, prevent the generation of new defects or grain boundaries during cooling, and effectively improve the stability and integrity of the single crystal structure.

[0089] For example, the first alloy is a NiTi alloy, which is annealed in a dual-temperature gradient zone of 1160 ℃ high temperature zone and 1100 ℃ low temperature zone. The temperature is raised synchronously at a rate of 10 ℃ / min, held at the temperature for 72 h and then cooled in the furnace. This can ensure the complete recrystallization and grain boundary migration of NiTi alloy while avoiding melting, abnormal grain growth and thermal stress cracks. It can also efficiently induce dominant orientation grains to penetrate the entire component, and finally obtain a NiTi alloy single crystal with high integrity and high stability.

[0090] Optionally, the temperature gradient annealing is performed under the protection of an inert gas and / or a reducing gas.

[0091] Optionally, the temperature gradient annealing is carried out under a 5 vol.% H2-Ar mixed gas protection environment.

[0092] This avoids oxidation of the first alloy, inhibits impurity intrusion, ensures compositional uniformity and grain boundary migration activity, and provides a clean and stable thermodynamic environment for directional growth of single crystals.

[0093] In some embodiments, during the heat preservation stage, the temperature difference between the third temperature and the fourth temperature is maintained at 40~80 ℃ to form a static temperature gradient of 0.25~0.35 ℃ / mm along the length direction of the first alloy single crystal.

[0094] Optionally, during the heat preservation stage, the temperature difference between the third temperature and the fourth temperature is maintained at 60 ℃ to form a static temperature gradient of 0.3 ℃ / mm along the length direction of the first alloy single crystal.

[0095] During the holding stage, the temperature difference between the third and fourth temperatures is maintained at 40~80 ℃, forming a static temperature gradient of 0.25~0.35 ℃ / mm along the length of the first alloy single crystal. This provides a continuous and stable growth driving force for dominant oriented grains (such as <001> orientation). This differs from the gradient-free environment of traditional homogeneous annealing, guiding grain boundaries to migrate directionally from the high-temperature region to the low-temperature region, avoiding polycrystalline formation caused by random grain boundary migration, and suppressing the nucleation of new grains and abnormal growth of brittle phases. Combined with slow heating and long-term holding, this ensures the uniformity of single crystallization in large-size components, ultimately improving the <001> orientation occupancy and density of the first alloy single crystal.

[0096] In some embodiments, the first alloy comprises a shape memory alloy.

[0097] The first alloy is a shape memory alloy. Using the aforementioned preparation method, shape memory alloy single crystals can be precisely controlled for preferred crystal orientation, such as the <001> orientation. This eliminates the hindering effect of grain boundaries on the nucleation, growth, and reverse phase transformation of martensite in polycrystalline shape memory alloys, effectively improving the shape memory recovery rate of the single-crystal alloy. Furthermore, it effectively controls the fluctuation range of the phase transformation temperature, enhancing the stability and reversibility of the shape memory function. Simultaneously, the single-crystal structure avoids stress concentration and crack initiation risks at grain boundaries. Combined with the inherent superelasticity of shape memory alloys, the prepared single crystals can effectively improve fatigue resistance and service life when subjected to cyclic loading or large elastic deformation. In addition, this preparation method can be synergistically combined with the composition control process of shape memory alloys. By adjusting the alloy element ratio, such as the percentage of Ni in a NiTi alloy, the phase transformation temperature requirements of different applications such as medical and industrial can be matched. The resulting shape memory alloy single crystals possess excellent biocompatibility and corrosion resistance.

[0098] In some embodiments, the first alloy includes any one of NiTi alloy, Cu-based shape memory alloy, Fe-based shape memory alloy, and TiNb-based β-type titanium alloy.

[0099] The first alloy can be a NiTi alloy, a Cu-based shape memory alloy, an Fe-based shape memory alloy, or a TiNb-based β-type titanium alloy. These alloys all possess reversible phase transformation properties and good additive manufacturing compatibility, enabling high-quality forming through vacuum environment and parameter control in electron beam fused deposition processes. The B2 alloy of NiTi... The B19' thermoelastic martensitic phase transformation endows it with shape memory effect and superelasticity. Besides NiTi alloys, the first alloy can also be a Cu-based shape memory alloy, an Fe-based shape memory alloy, or a TiNb-based β-type titanium alloy. Cu-based and Fe-based shape memory alloys are low-cost and can expand the application scenarios of single-crystal phase transformation alloys. TiNb-based β-type titanium alloys have excellent biocompatibility and can meet the needs of the medical field.

[0100] Optionally, the Ni atomic percentage of the NiTi alloy is 49.0~51.0 at.%. Optionally, the Ni atomic percentage of the NiTi alloy is 50.7 at.%.

[0101] The aforementioned structural and performance parameters of the NiTi alloy effectively match the requirements of electron beam fused deposition (EBD) and thermal cycling processes. Excessive or insufficient Ni content inhibits martensitic transformation, necessitating that the NiTi alloy maintain stable reversible phase transformation capability during the vacuum environment and thermal cycling of the EBD process. A Ni atomic percentage of 50.7 at.% in the NiTi alloy allows for a uniform B2 austenitic 601 matrix and trace amounts of residual B19′ martensite 602, while also reducing excessive Ni volatilization under high-energy electron beam bombardment. This ensures a precise and stable Ni / Ti atomic ratio, providing a high-quality microstructure with uniform composition for subsequent single-crystal growth.

[0102] Optionally, the austenitic transformation end temperature of the NiTi alloy is 25~35 ℃. Optionally, the austenitic transformation end temperature of the NiTi alloy is 29.8 ℃.

[0103] Optionally, the martensitic transformation end temperature of the NiTi alloy is -25 to -15 ℃. Optionally, the martensitic transformation end temperature of the NiTi alloy is -18.6 ℃.

[0104] The austenitic transformation temperature of NiTi alloys is 25–35 °C, close to room temperature, allowing NiTi alloys to exist primarily as the B2 austenitic phase at room temperature. The martensitic transformation temperature of NiTi alloys is -25 to -15 °C, enabling them to completely transform into the B19′ martensite 602 phase during cold immersion in 802 heat. This repeated phase transformation induces uniformly distributed dislocation defects, providing active sites for directional grain growth during temperature gradient annealing. This approach satisfies the premise of a reversible phase transformation where the first temperature of the thermal environment is higher than the austenitic transformation temperature of NiTi alloys, and the second temperature of the cold environment is lower than the martensitic transformation temperature of NiTi alloys. Furthermore, it exhibits good room-temperature superelasticity in subsequent applications, ensuring a moderate driving force for phase transformation during thermal cycling.

[0105] In some embodiments, step S110, before depositing the first alloy layer by layer using an electron beam fused wire deposition process, further includes: The first alloy was cleaned by immersion in acid and acetone, and then kept at a fifth temperature. The substrate was polished and the surface was wiped with acetone, and then kept at a fifth temperature.

[0106] Before depositing the first alloy layer by layer, the first alloy is subjected to pickling, acetone cleaning and heat treatment at the fifth temperature. This pretreatment can effectively remove surface impurities and oxide film from the first alloy, eliminate residual moisture and stress, and provide a clean and stable forming substrate for the electron beam fused wire deposition process. This effectively improves the forming quality, microstructure density and performance stability of the first intermediate.

[0107] Optionally, the first alloy is made of wire, and the diameter of the first alloy wire is 1.5~2.0 mm. Optionally, the diameter of the first alloy wire is 1.6 mm.

[0108] The diameter of the first alloy wire is selected from 1.5 to 2.0 mm, which can match the wire feeding stability and melt pool formation requirements of the electron beam fused wire deposition process. This avoids wire feeding jitter caused by excessively small diameter wires, and complete melting can be achieved without excessively high energy density. Combined with process parameters, it can ensure the continuity and density of layer-by-layer deposition. A diameter of 1.6 mm allows the melting rate of the first alloy wire to be better matched with the wire feeding speed.

[0109] Optionally, the fifth temperature is 50~70 ℃. Optionally, the fifth temperature is 60 ℃.

[0110] The fifth temperature setting, 50-70℃, avoids incomplete drying due to excessively low temperatures or oxidation of the wire surface caused by excessively high temperatures, effectively removing residual moisture from the surface of the first alloy wire and the substrate. Experiments have verified that a fifth temperature of 60℃ can achieve complete drying in a short time with the lowest energy consumption.

[0111] Optionally, the first alloy is cleaned by immersion in acid and acetone, and then heat-treated at a fifth temperature. After polishing the substrate and wiping the surface with acetone, and heat-treating at the fifth temperature, the process further includes: installing the treated first alloy wire into the wire feeding mechanism of the electron beam fused deposition apparatus, clamping the treated substrate onto the motion system in the vacuum chamber of the apparatus, and adjusting the positions of the wire feeding rod and the wire feeding nozzle so that the first alloy wire forms a first angle with the horizontal direction, and the tip of the first alloy wire maintains a first distance in height from the substrate surface.

[0112] Optionally, the first included angle is 35°~45°. Optionally, the first included angle is 40 degrees.

[0113] Optionally, the first spacing is 0.8 mm to 1.2 mm. Optionally, the first spacing is 1 mm.

[0114] After the first alloy wire is clamped, the position is calibrated, and the first included angle and the first spacing are matched to ensure that the droplets continuously flow into the molten pool through the liquid bridge transition, avoiding defects such as wire breakage and droplet splashing, and improving the continuity and density of the deposited layer.

[0115] Optionally, after each deposition-remelting cycle is completed, the stage height of the electron beam fused wire deposition equipment is lowered by a first distance, and a new first alloy layer is deposited.

[0116] Optionally, the first distance is 0.8 mm to 1.2 mm. Optionally, the first distance is 1 mm.

[0117] The table descends a certain distance each time, matching the melting rate of the wire and the thickness of the deposited layer. This improves the uniformity of the interlayer thickness, providing a uniform structural foundation for subsequent interlayer remelting, thermal cycling, and temperature gradient annealing, and reducing interlayer stress concentration.

[0118] This disclosure provides a phase change alloy single crystal, which is prepared using the method for preparing a phase change alloy single crystal without deformation treatment as described in any of the foregoing embodiments.

[0119] The phase transformation alloy single crystal provided in this embodiment is prepared using the aforementioned deformation-free preparation method, which eliminates the complex plastic deformation and subsequent annealing processes in traditional methods. This not only simplifies the preparation process, shortens the production cycle, and reduces process costs, but also avoids lattice defects and residual internal stress caused by deformation treatment. The prepared phase transformation alloy single crystal has precise and controllable orientation, no grain boundary interference within the crystal, and a uniform and stable phase transformation process. Its shape memory effect recovery rate and superelastic properties are significantly better than those of polycrystalline and deformed single-crystal materials, and its phase transformation temperature fluctuation range is narrow.

[0120] This disclosure provides a phase change alloy single crystal component, which is composed of a first alloy single crystal, and the first alloy single crystal is prepared by the preparation method described in any of the foregoing embodiments.

[0121] The phase change alloy single crystal component of this embodiment is composed of a first alloy single crystal as the core and is prepared by a synergistic process of electron beam filament deposition pretreatment, thermal cycling treatment and temperature gradient annealing. It has high density, high orientation consistency and low defect density, good shape memory effect and superelasticity, and strong structural stability.

[0122] In some embodiments, the first alloy single crystal has a <001> preferred orientation.

[0123] The first alloy single crystal exhibits a preferred <001> crystal orientation, which effectively optimizes its functional properties and mechanical stability. This orientation is similar to the B2 orientation of phase transformation alloys such as NiTi alloy. The B19′ reversible phase transition characteristics are well-suited to reduce the critical stress for stress-induced phase transitions, improving phase transition reversibility and response rate, and enhancing shape memory recovery rate and superelastic recovery ratio. Simultaneously, the <001> orientation reduces the hindrance of grain boundaries to phase transitions and lowers the stress threshold of Ti4Ni2O. x The aggregation of brittle phases at grain boundaries, combined with a temperature gradient annealing process, ensures the compositional uniformity and density of the first alloy single crystal. Furthermore, this preferred orientation enhances the mechanical properties of the first alloy single crystal along a specific direction, improving fatigue resistance and stability.

[0124] Optionally, the first alloy single crystal has a preferred orientation of <001> crystal direction, and the orientation deviation is less than or equal to 5°. This maximizes the consistency and stability of the functional performance of the phase change alloy single crystal component, and provides key structural guarantees for the excellent shape memory effect, superelasticity, and mechanical reliability of the first alloy single crystal.

[0125] In some embodiments, the single crystal grain size of the first alloy single crystal is on the order of millimeters.

[0126] The term "millimeter-scale" refers to the grain size of the first alloy single crystal being on the order of millimeters. The geometric dimensions of the grains are primarily measured in millimeters (mm), typically ranging from 0.1 mm to tens of millimeters. This embodiment of millimeter-scale single crystals balances fabrication feasibility with practical application requirements, avoiding the difficulty in processing micron-scale single crystals. The first alloy single crystal prepared using the aforementioned method achieves a grain size on the millimeter scale. Larger grains significantly reduce the number of grain boundaries, preventing stress concentration and defect aggregation at grain boundaries, reducing energy loss during phase transformation, and improving the uniformity and reversibility of phase transformation. Simultaneously, millimeter-scale grains reduce precipitation sites for brittle phases, minimizing obstacles to grain boundary migration.

[0127] This disclosure provides a terminal product, characterized in that it includes the aforementioned phase change alloy single crystal, or includes the aforementioned phase change alloy single crystal component.

[0128] The terminal product provided in this disclosure includes the aforementioned phase change alloy single crystal, or the aforementioned phase change alloy single crystal component, and therefore has all the beneficial effects of the aforementioned phase change alloy single crystal or phase change alloy single crystal component, which will not be repeated here.

[0129] The following specific embodiments illustrate the preparation method of the phase change alloy single crystal without deformation treatment and the phase change alloy single crystal of this disclosure, so as to more clearly illustrate the technical problems solved by this application, the technical solutions, and the beneficial effects. Obviously, the described embodiments are only some embodiments of this application, and not all embodiments. The following description of at least one exemplary embodiment is merely illustrative and is in no way intended to limit this application and its applications.

[0130] Where specific techniques or conditions are not specified in the examples, they shall be performed in accordance with the techniques or conditions described in the literature in this field or in accordance with the product instructions. Reagents or instruments whose manufacturers are not specified are all commercially available conventional products.

[0131] Example 1 Example 1 describes the pretreatment of the first alloy and the substrate. The first alloy is made of Ti-50.7at.%Ni alloy wire, and the substrate is made of TA1 pure titanium substrate.

[0132] (1.1) Pickling is performed on Ti-50.7at.%Ni alloy wire with a diameter of 1.6 mm to remove the surface oxide film, and then it is soaked in acetone to clean the surface of the alloy wire to remove oil and impurities. (1.2) Sand the TA1 pure titanium substrate with dimensions of 250 mm × 100 mm × 10 mm until the surface is smooth and clean, and then wipe the surface with acetone to remove oil and impurities from the substrate surface. (1.3) The Ti-50.7at.%Ni alloy wire and the TA1 pure titanium substrate were placed in a drying oven at 60 ℃ for drying treatment and kept at the temperature for 30 min to remove the moisture in the Ti-50.7at.%Ni alloy wire and the TA1 pure titanium substrate, and to eliminate the residual deformation in the Ti-50.7at.%Ni alloy wire, so that the Ti-50.7at.%Ni alloy wire is in a high-temperature austenitic state.

[0133] Example 2 In this embodiment 2, Ti-50.7at.%Ni alloy wire is deposited layer by layer and remelted interlayer using electron beam fused wire deposition.

[0134] (2.1) Clamping and positioning calibration of substrate and wire The treated Ti-50.7at.%Ni alloy wire is installed on the wire feeding mechanism of the electron beam fused deposition equipment. The treated TA1 pure titanium substrate is clamped on the motion system in the vacuum chamber of the equipment. The positions of the wire feeding rod and the wire feeding nozzle are adjusted so that the wire is at a 40-degree angle to the horizontal direction, and the tip of the wire is kept 1 mm away from the substrate surface in height to ensure that the droplets continuously flow into the molten pool through the liquid bridge transition.

[0135] (2.2) Deposition program settings for electron beam filament deposition equipment i. A beam current density of 25 mA was set to preheat the TA1 pure titanium substrate five times to promote the compatibility between the Ti-50.7at.%Ni deposition layer and the TA1 pure titanium substrate. ii. Set the following Ti-50.7at.%Ni alloy wire deposition process parameters: acceleration voltage is 60 kV, focusing current is 1000 mA, wire feed speed is 3000 mm / min, printing speed is 600 mm / min, beam current density is 45 mA (i.e., the first beam current density), and the moving trajectory is unidirectional scanning mode; iii. After each layer is deposited, a 30-second dwell time is set to allow the deposited layer to solidify. Then, the Ti-50.7at.%Ni alloy wire feed is stopped, and the current deposited layer is remelted at a beam current density of 20 mA (i.e., the second beam current density). iv. After each deposition-remelting cycle, the stage height is lowered by 1 mm before depositing a new Ti-50.7at.%Ni alloy layer.

[0136] (2.3) The vacuum level of the vacuum chamber of the electron beam fused deposition equipment reaches the required operating level (7×10). -2 When the concentration is below Pa, the electron beam filament deposition of Ti-50.7at.%Ni alloy is rapidly prepared according to the deposition procedure preset in step (2.2). The entire deposition-remelting process is as follows: Figure 6 As shown; the obtained Ti-50.7at.%Ni alloy bulk (i.e., the first intermediate) has dimensions of 220mm × 7.5mm × 25mm, exhibiting a strong cubic texture {001}〈001〉, as... Figure 9 and Figure 10 As shown.

[0137] In this embodiment, the testing methods used include X-ray diffraction analysis (XRD) and electron backscatter diffraction analysis (EBSD).

[0138] Figure 9 This is a crystallographic texture distribution diagram of the Ti-50.7at.%Ni alloy sample. Figure 9 As shown, the first intermediate, Ti-50.7at.%Ni alloy, exhibits a strong cubic texture {001}<001>, with a high degree of consistency in texture orientation along the deposition direction (vertical direction in the figure). The texture intensity is close to 100%, indicating that the {001} crystal planes of most grains are parallel to the deposition plane, and the <001> crystal orientations are aligned along the deposition direction. This lays the texture foundation for subsequent temperature gradient annealing to transform the crystals into <001> oriented single crystals. Figure 9There are subtle differences in color intensity in some areas. Combined with the scale (200 μm), it can be seen that the structure is a cubic textured aggregate of polycrystalline columnar crystals (not a single crystal). However, the strong texture characteristics can reduce the orientation competition for subsequent single crystallization.

[0139] Figure 10 This is the XRD pattern of a Ti-50.7at.%Ni alloy sample. The observed and calculated curves in the spectrum show a high degree of agreement. The refined results indicate that the first intermediate mainly consists of B2 austenite phase (90.8%) and B19' martensite phase (9.2%), with no Ti4Ni2O. x It contains brittle impurities with high phase purity. Fitting parameters (R0) wp =3.86%, R p =3.66%, CHI 2 =1.63) are all within a reasonable range, proving that the phase analysis results are accurate. Figure 10 The XRD pattern shows that the {001} crystal plane diffraction peak intensity of the B2 phase is relatively high, and... Figure 9 The strong cubic texture {001}〈001〉 forms a corresponding structure, which verifies the texture characteristics of the first intermediate from the perspective of phase composition, proving that although the first intermediate is polycrystalline, it has the characteristics of strong cubic texture and high phase purity.

[0140] Example 3 In this embodiment 3, the first intermediate Ti-50.7at.%Ni alloy block is subjected to thermal cycling treatment.

[0141] The Ti-50.7at.%Ni alloy block obtained in step (2.3) was cut into cubic samples with dimensions of 200 mm × 5 mm × 20 mm. After cleaning, the samples were immersed in a 90 ℃ constant temperature water bath (hot environment) for 3 min, then removed and immersed in liquid nitrogen (cold environment) for 3 min, and then removed and immersed in a 90 ℃ constant temperature water bath (hot environment) for 3 min. This "hot-cold-hot" cycle was repeated 50 times. The entire process is as follows: Figure 7 As shown; Figure 11This is a schematic diagram of the evolution of the internal microstructure of the Ti-50.7at.%Ni alloy sample during the thermal cycling process. Because the initial thermal environment temperature is much higher than the austenitic transformation end temperature of the Ti-50.7at.%Ni alloy (measured Af = 29.8 ℃), it exhibits a fully austenitic microstructure. However, the subsequent cold environment temperature is much lower than the martensitic transformation end temperature of the Ti-50.7at.%Ni alloy (measured Mf = -18.6 ℃), which promotes the full occurrence of the martensitic transformation (austenite-martensite). When placed back into the thermal environment, the reverse martensitic transformation (martensite-austenite) will occur. The reversible transformation corresponding to this cycle process will induce plastic deformation within the grains, ultimately forming a high-density dislocation substructure inside the Ti-50.7at.%Ni alloy sample (i.e., the second intermediate).

[0142] Example 4 Example 4 involves subjecting the second intermediate Ti-50.7at.%Ni alloy sample to temperature gradient annealing.

[0143] Combination Figure 8 As shown, the Ti-50.7at.%Ni alloy sample treated in step (3) was placed in a dual-temperature zone tube furnace, evacuated and 5% H2-Ar mixed gas was introduced. The target temperature of the high-temperature zone was set to 1160 ℃, and the target temperature of the low-temperature zone was set to 1100 ℃. The two temperature zones were simultaneously raised from room temperature to the target temperature at a rate of 10 ℃ / min. After stabilization, the temperature was held for 72 hours, and finally cooled to room temperature with the furnace. According to calculations, a static temperature gradient of 0.3 ℃ / mm was formed along the length direction (~200 mm) of the Ti-50.7at.%Ni alloy sample during the holding stage. Combined with the high-density dislocation substructure formed in step (2.3), abnormal grain growth can be used to form seed crystals and drive their directional expansion, ultimately transforming the cubic textured columnar crystals into <001> oriented single crystals. Figure 12 As shown, the left-hand diagram is the corresponding {001} pole figure. The diffraction spots in the pole figure are highly concentrated, exhibiting a typical single-point dense distribution characteristic, with no dispersed diffraction signals. Combined with pole figure calibration rules, this distribution indicates that the {001} crystal planes of all grains in the Ti-50.7at.%Ni alloy sample are parallel to the same reference plane, proving the formation of a single <001> preferred orientation with an orientation deviation ≤5°. The right-hand diagram is the standard inverse pole figure of B2 austenite. The color distribution in the inverse pole figure is uniform, showing only a single dominant orientation color region, with no color impurities corresponding to other orientations. Combined with the standard inverse pole figure scale, this dominant orientation matches the <001> crystal orientation, confirming that the seed crystal formed by abnormal grain growth in the high-temperature region has directionally absorbed all non-<001> oriented grains along the temperature gradient direction, achieving a complete transformation from cubic textured columnar crystals to a single <001> oriented single crystal.

[0144] In this embodiment, the testing method used includes electron backscatter diffraction analysis (EBSD).

[0145] Comparative Example 1 The main difference between Comparative Example 1 and Examples 1 to 4 is that the first alloy uses Ti-50.0 at.%Ni alloy wire, the deposition process uses only single electron beam wire deposition, without layer-by-layer deposition-interlayer remelting cycle operation, and without the thermal cycling treatment of Example 3 and the temperature gradient annealing of Example 4; it is formed only by natural cooling during the deposition process. The specific process of Comparative Example 1 is as follows: (1) Pickling the Ti-50.0at.%Ni alloy wire with a diameter of 1.6 mm to remove the surface oxide film, and then immersing it in acetone to ensure that there is no oil or impurities on the surface of the Ti-50.0at.%Ni alloy wire; (2) Sand the TA11 pure titanium substrate until the surface is smooth and clean, and then wipe the surface with acetone to remove oil and impurities from the substrate surface. (3) The Ti-50.0at.%Ni alloy wire and the TA11 pure titanium substrate were placed in a drying oven at 60°C for drying treatment and kept at the temperature for 30 min to remove the moisture in the Ti-50.0at.%Ni alloy wire and the TA11 pure titanium substrate, and to eliminate the residual deformation in the Ti-50.0at.%Ni alloy wire, so as to ensure that the Ti-50.0at.%Ni alloy wire is in the high-temperature austenitic state; (4) The treated Ti-50.0at.%Ni alloy wire is installed on the wire feeding mechanism of the electron beam fused wire deposition equipment, and the treated pure titanium substrate is clamped on the motion system in the vacuum chamber of the equipment. (5) The vacuum level of the vacuum chamber of the electron beam fused deposition equipment reaches the required operating level (7×10). -2 When the speed is Pa), the following processing parameters are set: acceleration voltage U is 60 kV, focusing current If is 1000 mA, wire feed speed Vfeed is 3000 mm / min, printing speed Vprint is 600 mm / min, and beam current density Ib is 45 mA. (6) Edit the printing program under the process parameters in step (5) in the automatic operation window of the electron beam fused wire deposition equipment according to the reciprocating scanning mode; (7) Following the printing procedure in step (6), complete the rapid preparation of NiTi alloy by electron beam filament deposition to obtain gradient NiTi alloy manufactured by electron beam filament additive manufacturing.

[0146] The phase spectrum of the NiTi alloy prepared in Comparative Example 1 is as follows: Figure 13 As shown. From Figure 13The phase spectrum shown reveals that, along the deposition direction, the NiTi alloy manufactured by electron beam filament additive manufacturing can be divided into four regions: transition zone, bottom, middle, and top. The phase composition of each region differs significantly. In the transition zone, the Ti2Ni matrix is ​​dominant, with small amounts of austenitic TiNi and martensitic TiNi solid solutions forming. As the deposition height increases, austenitic TiNi solid solutions become the main phase, with small amounts of martensitic TiNi solid solutions and Ti2Ni intermetallic compounds precipitating. Furthermore... Figure 14 The solidification orientation of the NiTi alloy fabricated by electron beam filament additive manufacturing in Comparative Example 1 is presented. It can be observed that in addition to coarse columnar crystals with (001) orientation, there are also fine, uniform, isotropic equiaxed crystals between layers. Therefore, it can be concluded that the TiNi shape memory alloy fabricated by electron beam filament additive manufacturing in Comparative Example 1 exhibits a gradient change in phase and microstructure along the deposition height, which is a typical gradient NiTi alloy.

Claims

1. A method for producing a single crystal of a phase change alloy without a deformation process, characterized by, include: An electron beam fused wire deposition process is used to deposit the first alloy layer by layer and then remelt it between layers to obtain a first intermediate; wherein the first alloy has reversible martensitic phase transformation characteristics. The first intermediate is subjected to thermal cycling to obtain the second intermediate; The second intermediate was subjected to temperature gradient annealing to obtain the first alloy single crystal.

2. The production method according to claim 1, characterized by, An electron beam fused wire deposition process is used to deposit the first alloy layer by layer, followed by interlayer remelting to obtain a first intermediate, comprising: An electron beam is used to deposit a first alloy on a substrate in a unidirectional scanning mode at a first beam current density to form a phase change alloy deposition layer. An electron beam with a second beam current density is used to remelt the current phase change alloy deposit layer, forming a molten pool on the surface of the current phase change alloy deposit layer. A first alloy layer is deposited on the surface of the molten pool; A deposition-remelting cycle is performed to obtain the first intermediate; Optionally, the second beam current density is 20~25 mA; or, the second beam current density is 20 mA; optionally, the substrate is a TA1 pure titanium substrate.

3. The preparation method according to claim 2, characterized in that, The second beam density during interlayer remelting is lower than the first beam density during the deposition stage.

4. The preparation method according to claim 1, characterized in that, The first intermediate is subjected to thermal cycling to obtain the second intermediate, which includes: The first intermediate is alternately immersed in a hot environment at a first temperature and a cold environment at a second temperature to form a two-way cycle of hot to cold and cold to hot. The first temperature is higher than the austenitic phase transformation end temperature of the first alloy, and the second temperature is lower than the martensitic phase transformation end temperature of the first alloy.

5. The preparation method according to claim 4, characterized in that, The thermal environment adopts a water bath environment; and / or The first temperature is 40~100 ℃; or, the first temperature is 90 ℃; and / or Each immersion in a hot environment takes 2-5 minutes; or, each immersion in a hot environment takes 3 minutes; and / or The cold environment is provided by liquid nitrogen; and / or The second temperature is -196 to -30 ℃; or, the second temperature is -196 ℃; and / or Each cold immersion session should last 2–5 minutes; or, each cold immersion session should last 3 minutes; and / or The number of alternating soaking cycles is 40-80 times; or, the number of alternating soaking cycles is 50 times.

6. The preparation method according to claim 1, characterized in that, The second intermediate is subjected to temperature gradient annealing to obtain the first alloy single crystal, including: The second intermediate is placed in a dual-temperature zone device, the target temperature of the high-temperature zone is set as the third temperature, the target temperature of the low-temperature zone is set as the fourth temperature, the high-temperature zone and the low-temperature zone are heated to the target temperature at the first rate, then kept at the temperature for the first time, and then cooled to room temperature. Optionally, the third temperature is 1120~1200 ℃; optionally, the third temperature is 1160 ℃; optionally, the fourth temperature is 1060~1140 ℃; optionally, the fourth temperature is 1100 ℃; optionally, the first rate is 5~15 ℃ / min; optionally, the first rate is 10 ℃ / min; optionally, the first duration is 48~96 h; optionally, the first duration is 72 h.

7. The preparation method according to claim 6, characterized in that, During the heat preservation stage, the temperature difference between the third and fourth temperatures is maintained at 40~80 ℃, so as to form a static temperature gradient of 0.25~0.35 ℃ / mm along the length direction of the first alloy single crystal; or, During the heat preservation stage, the temperature difference between the third and fourth temperatures is maintained at 60 ℃ to form a static temperature gradient of 0.3 ℃ / mm along the length of the first alloy single crystal.

8. The preparation method according to any one of claims 1 to 7, characterized in that, The first alloy includes a shape memory alloy; Optionally, the first alloy includes any one of NiTi alloy, Cu-based shape memory alloy, Fe-based shape memory alloy, and TiNb-based β-type titanium alloy; Optionally, the Ni atomic percentage of the NiTi alloy is 49.0~51.0 at.%; or, the Ni atomic percentage of the NiTi alloy is 50.7 at.%. Optionally, the austenitic transformation end temperature of the NiTi alloy is 25~35 ℃; or, the austenitic transformation end temperature of the NiTi alloy is 29.8 ℃. Optionally, the martensitic transformation end temperature of the NiTi alloy is -25 to -15 ℃; or, the martensitic transformation end temperature of the NiTi alloy is -18.6 ℃.

9. A phase transformation alloy single crystal, characterized in that, It is prepared by the method for preparing phase transformation alloy single crystals without deformation treatment as described in any one of claims 1 to 8.

10. A phase change alloy single crystal component, characterized in that, The phase change alloy single crystal component is composed of a first alloy single crystal, which is prepared by the preparation method of phase change alloy single crystal without deformation treatment as described in any one of claims 1 to 8.

11. The phase change alloy single crystal component according to claim 10, characterized in that, The first alloy single crystal has a preferred orientation of <001>; and / or The single crystal grain size of the first alloy single crystal is on the millimeter scale.

12. A terminal product, characterized in that, It includes the phase change alloy single crystal as described in claim 9, or the phase change alloy single crystal component as described in claim 10 or 11.

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