Heterostructure reinforced wide-temperature-range Ni-Mo-Al alloy and preparation method thereof
By introducing Al into Ni-Mo alloys and constructing heterostructures, the problem of mid-temperature brittleness in high-molybdenum-content Ni-Mo alloys was solved, achieving a balance between strength and plasticity over a wide temperature range, resulting in good mechanical properties and industrialization potential.
Patent Information
- Authority / Receiving Office
- CN · China
- Patent Type
- Applications(China)
- Current Assignee / Owner
- HUNAN UNIV
- Filing Date
- 2026-02-03
- Publication Date
- 2026-05-01
AI Technical Summary
High molybdenum-content Ni-Mo binary alloys are prone to forming brittle phases in the mid-temperature region, leading to loss of plasticity. Existing technologies make it difficult to effectively control the ordered behavior through alloying elements and achieve a balance between strength and plasticity over a wide temperature range.
By introducing Al into Ni-Mo alloys and combining cold working deformation with annealing heat treatment, a partially recrystallized heterogeneous structure is constructed, forming a heterogeneous structure of non-recrystallized and recrystallized regions. This regulates the ordered behavior to alleviate brittleness and achieves a balance between strength and plasticity.
Achieving a balance between strength and plasticity at extreme brittle temperatures, Ni-Mo-Al alloys exhibit high yield strength and good elongation in the mid-temperature range, with a wide temperature range where strength and plasticity are well-matched. Moreover, they are low in cost and have simple processes that are easy to industrialize.
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Abstract
Description
A wide-temperature-range Ni-Mo-Al alloy strengthened by heterostructure and its preparation method Technical Field
[0001] This invention relates to the field of alloy materials technology, specifically to a wide-temperature-range Ni-Mo-Al alloy strengthened by a heterostructure and its preparation method. Background Technology
[0002] With the increasing demands for adaptability to extreme environments in fields such as advanced aerospace and cryogenic engineering, the development of structural materials that can maintain high plasticity and avoid catastrophic failure in a wide temperature range from ultra-low temperatures (such as -196℃) to medium and high temperatures (600-800℃) has become an urgent need.
[0003] Among numerous material systems, Ni-Mo binary alloys with high molybdenum (Mo) content have attracted much attention due to their excellent strength-plasticity balance at low and room temperatures. However, high Mo content presents the problem of ordered transformation in the mid-temperature region, often leading to catastrophic failure when these alloys are used at specific temperatures. 80 Mo 20 It is one of the representative alloys, which forms a large amount of brittle D1 at 750℃. a This phase causes it to completely lose its plasticity, becoming the fundamental bottleneck restricting its wide-temperature-range application.
[0004] In existing technologies, researchers often introduce alloying elements to regulate the ordering behavior of Ni-Mo binary alloys or alter the precipitation pathway of brittle ordered phases. Specifically, to regulate the ordering behavior of Ni-Mo binary alloys, researchers often introduce alloying elements (such as Al) to change the phase transformation pathway. It is worth noting that while existing studies have systematically explored the effects of Al addition (2–9 at.%) on Ni... 80 Mo 20 The influence of alloy microstructure (Shen, E., Brooks, CR, & Kenik, EA, Effect of 2, 5, 7, and 9 At.% Al on the solution heat-treated (1280°C, 50 hours, water quenched) structure of Ni-20at.%Mo alloy, Materials Characterization, 1999, 43: 57–68.) was studied, but the research conditions mainly focused on the quenched state after high-temperature solution / aging treatment (1250°C / 1280°C), and the metastable DO observed was... 22The formation of the Ni3Mo and even Ni2Mo phases reflects the microstructure frozen during high-temperature quenching. However, the phase evolution and microstructure stability of alloys in practical mid-temperature applications often differ significantly from the high-temperature quenched state. Furthermore, in aging studies closer to mid-temperature conditions (Kulkarni, UD, &Dey, GK, Ordering and topologically close acked-phase precipitation in aNi-25at.% Mo-5at.% Al alloy, Acta Materialia, 2004, 52: 2711–2720.), it was found that Ni... 70 Mo 25 When Al5 alloy is aged at 875K, it evolves from a short-range ordered state with diffraction characteristics of {1½0}, {100}, and {110} into D0. 22 Ni3Mo is formed, but with further increases in temperature or time, D0 will form. 22 A mixed microstructure consisting of two phases, Ni3Mo and Pt2Mo, coexisting. Further theoretical simulation studies (Kulkarni, UD, Montecarlo simulation of ordering transformations in Ni–Mo-based alloys, ActaMaterialia, 2004, 52: 2721–2732.) revealed the mechanism by which Al affects phase stability at the atomic scale: the addition of Al causes the short-range ordered state to change from D0 in the binary alloy. 22 Competition with the N3M structure controlled the shift to a significant bias towards D0. 22 This represents an early, ordered evolutionary state. This facilitates the development of D0 by Al. 22 The formation provides a theoretical explanation.
[0005] However, a systematic understanding of the microstructure evolution, phase stability, and final mechanical properties of alloys with Al addition under prolonged exposure in the mid-temperature range remains lacking. Therefore, although Al addition provides the possibility of controlling the ordering pathway, how to address the mid-temperature brittleness problem that high-molybdenum (Mo) content Ni-Mo binary alloys may face in the mid-temperature region, and how to achieve a synergistic effect of microstructure stability and plasticity by suppressing brittle phases through compositional design and process control, remains a weak link in current research.
[0006] To synergistically enhance the strength and ductility of materials, constructing heterogeneous structures (such as bimodal grains, layered or partially recrystallized structures) has become a cutting-edge alloy design strategy. This strategy achieves a balance between material strength and ductility through the synergistic deformation of soft and hard regions during deformation, inducing strengthening and hardening through heterogeneous deformation. Among these, partially recrystallized structures, as a typical heterogeneous structure, have shown potential in improving the mid-temperature ductility of certain materials.
[0007] However, directly applying the aforementioned heterostructure strategy to solve the mid-temperature brittleness problem faced by high-molybdenum-content Ni-Mo alloys still presents fundamental challenges: First, in the sensitive temperature range, the tendency of the matrix to transform into a brittle ordered phase is the fundamental origin of its brittleness, and microstructure adjustments alone are insufficient to overcome this essential problem; second, under high-temperature conditions, the stability of the hard regions (non-recrystallized regions) in the heterostructure is tested, as dynamic processes such as recovery and recrystallization that may occur during deformation can affect their effectiveness as a hard phase. Therefore, in high-molybdenum-content Ni-Mo alloys, how to combine alloy design with microstructure control through processes suitable for bulk materials, while controlling the ordered behavior to alleviate brittleness, and precisely controlling the microscopic evolution of the hard regions at high temperatures to maintain their role as hard regions, thereby alleviating or overcoming brittleness problems at specific temperatures and achieving optimization of strong plasticity over a wide temperature range, is a technical challenge that urgently needs to be overcome. Summary of the Invention
[0008] The purpose of this invention is to provide a wide-temperature-range Ni-Mo-Al alloy strengthened by a heterostructure and its preparation method. By controlling the Al element content and combining cold working deformation and annealing heat treatment processes, a partially recrystallized heterostructure is constructed in the Ni-Mo-Al alloy, thereby achieving a balance between strength and plasticity at extreme brittle temperatures and a good match between strength and plasticity over a wide temperature range.
[0009] The first aspect of this invention is to provide a wide-temperature-range Ni-Mo-Al alloy strengthened by a heterostructure, the technical solution of which is:
[0010] A wide-temperature-range Ni-Mo-Al alloy strengthened by a heterostructure has the following elemental percentage composition expression: (Ni 80+y Mo 20-y ) 100-x Al x Where 2≤x≤8, -5≤y≤2; the Ni-Mo-Al alloy has a face-centered cubic FCC matrix, and the face-centered cubic FCC matrix contains a heterostructure, the heterostructure is composed of recrystallized regions and non-recrystallized regions, wherein the volume fraction of the recrystallized regions is greater than or equal to 50%, and the volume fraction of the non-recrystallized regions is greater than or equal to 20% and less than 50%.
[0011] Preferably, the volume fraction of the non-recrystallized region is 20-40%.
[0012] Preferably, in the elemental expression of the Ni-Mo-Al alloy, 3≤x≤5 and -4≤y≤0.
[0013] Furthermore, the matrix contains a diffusely distributed σ phase, which is a Mo-rich phase.
[0014] Furthermore, the non-recrystallized region is layered and mainly consists of strip-shaped substructures divided by dislocation entanglements and deformation bands.
[0015] The second aspect of this invention is to provide a method for preparing a wide-temperature-range Ni-Mo-Al alloy strengthened by a heterostructure, the technical solution of which is:
[0016] A method for preparing a wide-temperature-range Ni-Mo-Al alloy strengthened by a heterostructure includes the following steps:
[0017] Step S1, vacuum melting and vacuum casting: Ni, Mo and Al metal raw materials are vacuum melted and vacuum cast according to the alloy composition ratio to obtain alloy ingots;
[0018] Step S2, homogenization process;
[0019] Step S3, cold deformation treatment: multi-pass rolling is carried out at room temperature, wherein the rolling amount per pass is 2-6% and the total rolling amount is 60-90%;
[0020] Step S4, Annealing heat treatment: Annealing heat treatment is performed in an atmospheric atmosphere to obtain a wide-temperature-range Ni-Mo-Al alloy strengthened with a heterostructure; the annealing heat treatment temperature is 870℃-1150℃, and the holding time is 1.5-20min. The annealing temperature can be 870℃, 900℃, 950℃, 1000℃, 1050℃, or 1100℃, or other values within this range. The temperature and time of the heat treatment can be adjusted according to the alloy material composition, especially the amount of Al added, with the aim of obtaining a heterostructure within the face-centered cubic (FCC) matrix structure.
[0021] Furthermore, step S1 also includes a surface cleaning pretreatment step for Ni, Mo, and Al metal raw materials before vacuum melting, wherein the surface cleaning pretreatment is performed by ultrasonic cleaning with anhydrous ethanol, followed by drying.
[0022] During smelting, Ni, Mo, and Al metal raw materials are stacked according to their melting points, with the order from top to bottom being Mo, Ni, and Al metal raw materials;
[0023] The melting process is carried out under an argon atmosphere by electric arc melting with an arc current of 300-380A. Each melting session lasts for 60-120 seconds. After each melting session, the sample is cooled for 5-10 minutes, then flipped over and remelted. The number of remelting sessions is 5-10.
[0024] Furthermore, in step S2, the homogenization process is carried out in an inert gas atmosphere at 1150℃-1250℃ for 2-24 hours.
[0025] Furthermore, in step S3, the total rolling volume is 80-90%.
[0026] Furthermore, in step S2, after homogenization, water quenching is used for cooling; in step S4, after annealing heat treatment, water quenching is used for cooling.
[0027] Compared with the prior art, the wide-temperature-range Ni-Mo-Al alloy with heterostructure reinforcement and its preparation method provided by the present invention have the following advantages:
[0028] I. The present invention relates to a wide-temperature-range Ni-Mo-Al alloy strengthened by a heterostructure. By introducing the alloying element Al into the Ni-Mo alloy system, the ordered behavior of high-molybdenum-content Ni-Mo alloys in the mid-temperature range is slowed down or altered. Simultaneously, through the design of the heterostructure, a heterogeneous microstructure of non-recrystallized regions (hard regions) and recrystallized regions (soft regions) is constructed within the alloy material. Heterogeneous deformation-induced strengthening and hardening are utilized to achieve a balance between strength and ductility at a temperature traditionally prone to brittleness (750℃). Testing shows that at a mid-temperature (750℃), this alloy can possess a high yield strength of 690 MPa and a good total elongation of 29%.
[0029] II. The Ni-Mo-Al alloy of the present invention, strengthened by a heterostructure, exhibits a wide-temperature-range Ni-Mo alloy with a yield strength of approximately 1200 MPa, a tensile strength of approximately 1700 MPa, and a total elongation of approximately 40% at liquid nitrogen temperature (-196°C). At room temperature (20°C-25°C) and 500°C, its yield strength reaches approximately 1000 MPa and 900 MPa, its tensile strength reaches approximately 1350 MPa and 1300 MPa, and its total elongation reaches approximately 35% and 27%, respectively. Therefore, the Ni-Mo-Al alloy of the present invention achieves a good balance between strength and plasticity over a wide temperature range.
[0030] Third, the wide-temperature-range Ni-Mo-Al alloy strengthened by heterostructure of the present invention uses conventional pure metals as raw materials, and the added Al is a common lightweight alloying element, which avoids the use of expensive or scarce alloying elements and has the advantage of low cost.
[0031] Fourth, the preparation method of the wide-temperature-range Ni-Mo-Al alloy strengthened by heterostructure of the present invention has the advantages of simple process, short process and safety, which is conducive to realizing industrial production. Attached Figure Description
[0032] To more clearly illustrate the technical solutions in the embodiments of the present invention or the prior art, the drawings used in the description of the embodiments or the prior art will be briefly introduced below. Obviously, the drawings described below are some embodiments of the present invention. For those skilled in the art, other drawings can be obtained based on these drawings without creative effort.
[0033] Figure 1 shows the (Ni) in Example 1. 80 Mo 20 ) 96 Tensile stress-strain curves of Al4 alloy annealed at 900℃ for 2 min after cold rolling (Al) at -196℃, 25℃, 500℃, and 750℃.
[0034] Figure 2 shows the (Ni) in Example 1. 80 Mo 20 ) 96 XRD pattern of the initial microstructure of Al4 alloy after cold rolling and annealing at 900℃ for 2 min (A1).
[0035] Figure 3 shows the (Ni) in Example 1. 80 Mo 20 ) 96 SEM images of the initial microstructure of Al4 alloy after cold rolling and annealing at 900℃ for 2 min (A1) at different multiples along the TD direction.
[0036] Figure 4 shows the (Ni) in Example 1. 80 Mo 20 ) 96 SEM images of the initial microstructure of Al4 alloy after cold rolling and annealing at 900℃ for 2 min (Al) at different multiples along the RD direction.
[0037] Figure 5 shows the (Ni) in Example 1. 80 Mo 20 ) 96 SEM images of the initial microstructure of Al4 alloy after cold rolling and annealing at 900℃ for 2 min (A1) at different multiples along the ND direction.
[0038] Figure 6 shows the (Ni) in Example 1. 80 Mo 20 ) 96 The initial microstructure of Al4 alloy annealed at 900℃ for 2 min after cold rolling (Al) is shown in
[110] . FCCSelected area electron diffraction pattern of zone axis.
[0039] Figure 7 shows the (Ni) in Example 1. 80 Mo 20 ) 96 The alloy obtained by annealing Al4 alloy at 900℃ for 2 min after cold rolling (A1) and then performing SEM images of the sample with ~2.2% strain interruption at 750℃ along the TD direction at different magnifications.
[0040] Figure 8 shows the (Ni) in Example 1. 80 Mo 20 ) 96 The Al4 alloy was annealed at 900°C for 2 min after cold rolling (Al). A sample with a ~2.2% strain interruption was obtained at 750°C.
[110] FCC Zone axis,
[001] FCC Zone axis,
[112] FCC Selected area electron diffraction pattern of zone axis.
[0041] Figure 9 shows the (Ni) in Example 1. 80 Mo 20 ) 96 Fracture morphology and lateral microstructure near the fracture surface of an Al4 alloy obtained by annealing at 900℃ for 2 min after cold rolling (Al) at 750℃ under tensile fracture conditions at different ratios.
[0042] Figure 10 shows (Ni) 80 Mo 20 ) 96 SEM and XRD patterns of the initial microstructure of Al4 alloy after cold rolling and annealing at 900℃ for 1 h.
[0043] Figure 11 shows the (Ni) in Example 2. 80 Mo 20 ) 96 Tensile stress-strain curves of Al4 alloy obtained by annealing at 870℃ for 2 min after cold rolling (A2) at 25℃, 500℃, and 750℃.
[0044] Figure 12 shows the (Ni) in Comparative Example 1. 80 Mo 20 ) 96 The tensile stress-strain curves of the Al4 alloy after cold rolling and annealing at 1100℃ for 5 min (A3) at -196℃, 25℃, 500℃ and 750℃.
[0045] Figure 13 shows the (Ni) in Comparative Example 1. 80 Mo 20 ) 96XRD pattern of the initial microstructure of Al4 alloy after cold rolling and annealing at 1100℃ for 5 min (A3).
[0046] Figure 14 shows the (Ni) in Comparative Example 1. 80 Mo 20 ) 96 SEM image of the initial microstructure of Al4 alloy after cold rolling and annealing at 1100℃ for 5 min (A3).
[0047] Figure 15 shows the (Ni) in Comparative Example 1. 80 Mo 20 ) 96 The Al4 alloy obtained by cold rolling and annealing at 1100℃ for 5 min (A3) showed tensile fracture at 750℃ in
[110] . FCC Zone axis,
[001] FCC Selected area electron diffraction pattern of zone axis.
[0048] Figure 16 shows the (Ni) in Comparative Example 1. 80 Mo 20 ) 96 Fracture morphology and lateral microstructure near the fracture surface of an Al4 alloy obtained by annealing at 1100℃ for 5 min after cold rolling (A3) at 750℃.
[0049] Figure 17 shows Ni in Comparative Example 2. 80 Mo 20 The tensile stress-strain curves of the alloy obtained by annealing at 900℃ for 4 min after cold rolling (A4) at -196℃, 25℃, 500℃ and 750℃.
[0050] Figure 18 shows Ni in Comparative Example 2. 80 Mo 20 The XRD pattern of the initial microstructure of the alloy after cold rolling and annealing at 900℃ for 4 min (A4).
[0051] Figure 19 shows Ni in Comparative Example 2. 80 Mo 20 SEM image of the initial microstructure obtained after annealing the alloy at 900℃ for 4 min after cold rolling (A4).
[0052] Figure 20 shows Ni in Comparative Example 3. 80 Mo 20 The tensile stress-strain curves of the alloy obtained by annealing at 1100℃ for 5 min after cold rolling (A5) at -196℃, 25℃, 500℃ and 750℃.
[0053] Figure 21 shows Ni in Comparative Example 3. 80 Mo 20The XRD pattern of the initial microstructure obtained after annealing the alloy at 1100℃ for 5 min after cold rolling (A5).
[0054] Figure 22 shows Ni in Comparative Example 3. 80 Mo 20 SEM image of the initial microstructure obtained after annealing the alloy at 1100℃ for 5 min after cold rolling (A5).
[0055] Figure 23 shows the (Ni) in Example 1. 80 Mo 20 ) 96 The tensile mechanical properties of Al4 alloy obtained by annealing at 900℃ for 2 min after cold rolling (A1) are compared with those of nickel-based alloys, high-entropy alloys and other alloys at around 750℃. Detailed Implementation
[0056] To enable those skilled in the art to better understand the technical solutions in the embodiments of the present invention, and to make the above-mentioned objectives, features and advantages of the present invention more apparent and understandable, the specific embodiments of the present invention will be further described below.
[0057] The endpoints and any values of the ranges disclosed herein are not limited to the precise ranges or values, and these ranges or values should be understood to include values close to these ranges or values. For numerical ranges, the endpoint values of the various ranges, the endpoint values of the various ranges and individual point values, and individual point values can be combined with each other to obtain one or more new numerical ranges, which should be considered as specifically disclosed herein.
[0058] In this invention, when "room temperature" or "normal temperature" is used, the temperature can be 15-25℃.
[0059] Unless otherwise specified, all raw materials, reagents, instruments and equipment used in this invention can be purchased from the market or prepared by current methods.
[0060] Example 1
[0061] A wide-temperature range (Ni) reinforced with heterostructure 80 Mo 20 ) 96 The preparation method of Al4 alloy includes the following steps:
[0062] Step S1, vacuum melting and vacuum casting:
[0063] Ni, Mo, and Al metals with a purity ≥99.9% were selected as raw materials. After ultrasonic cleaning and drying with anhydrous ethanol, they were weighed according to the nominal alloy composition converted to a mass ratio (total weight 100g). The metal raw materials were placed in the copper mold crucible of the vacuum arc melting furnace in the order of Mo, Ni, and Al from top to bottom. After the furnace was closed, repeated vacuuming and argon purging were performed: first, the pressure was evacuated to 2×10⁻³ Pa, then 0.5 atm of high-purity argon was introduced, and this process was repeated 3 times. Finally, melting was carried out under 0.5 atm of argon. Before melting, the titanium ingot was melted at a current of 150 A for 90s to absorb residual oxygen in the furnace; then the current was adjusted to 60 A, the tungsten electrode head was moved above the alloy material, and the current was gradually increased to 350 A and held for 90s to fully melt and homogenize the alloy. Then the current was slowly reduced to 60 A and the furnace was shut off. This melting process was repeated 6 times. After each melting, the ingot was cooled for 5 minutes and turned over to ensure uniform composition. Finally, the alloy is transferred to the copper mold for suction casting, and the titanium ingot is melted again for 120 seconds (150A). Then, the alloy is quickly melted and suction cast with a current of 380A. After the alloy melts and flows into the copper mold, the current is immediately turned off to obtain an alloy cuboid ingot with a size of approximately 16×16×45mm.
[0064] Step S2, homogenization process:
[0065] The ingot was placed in a quartz glass tube with a wall thickness of 1 mm and a diameter of 25 mm, and a quartz glass column was placed at the top. The quartz tube was evacuated (to 2 × 10⁻¹ Pa) using a tube sealing machine and filled with 0.5 atm high-purity argon gas. After repeating the gas purging 10 times, the tube opening was sealed with an oxyhydrogen flame to form a complete seal. The sealed sample was placed in a muffle furnace and kept at 1200℃ for 24 h for homogenization treatment. After the treatment, the sample was quickly removed, the quartz tube was broken, and the sample was quenched in water for rapid cooling.
[0066] Step S3, cold deformation treatment:
[0067] The oxide scale on the homogenized alloy surface was removed by sanding with 180-grit sandpaper, and then multi-pass rolling was performed at room temperature, with the reduction per pass controlled at 6% and the total rolling deformation reaching 90%.
[0068] Step S4, Annealing heat treatment:
[0069] Dog-bone shaped tensile test specimens were cut along the rolled sheet, ensuring that the parallel sections of the specimens were parallel to the rolling direction. After grinding to remove cutting marks, the specimens were placed in a muffle furnace for annealing. Annealing was performed in air at 900°C for 2 minutes, followed immediately by water quenching.
[0070] The heterostructure-reinforced wide-temperature-range (Ni) alloy prepared in this embodiment 80 Mo 20 ) 96The Al4 alloy was named A1, and its initial microstructure was characterized. Its mechanical properties were tested at liquid nitrogen (-196℃), room temperature, 500℃, and 750℃. Specifically, the microstructure of the ~2.2% strain tensile interruption specimen and the tensile fracture specimen at 750℃ were characterized, and the fracture morphology of the tensile fracture surface was observed.
[0071] From the example (Ni) 80 Mo 20 ) 96 The engineering stress-strain curves of Al4 alloy obtained after tensile tests at different temperatures (as shown in Figure 1) show that at -196℃, its yield strength is approximately 1200 MPa, tensile strength is approximately 1700 MPa, and total elongation is approximately 40%; at room temperature, its yield strength is approximately 1030 MPa, tensile strength is approximately 1350 MPa, and total elongation is approximately 35%; at 500℃, its yield strength is approximately 900 MPa, tensile strength is approximately 1300 MPa, and total elongation is approximately 27%; and at 750℃, its yield strength is approximately 690 MPa, tensile strength is approximately 760 MPa, and total elongation is approximately 29%. X-ray diffraction analysis of the initial microstructure revealed a face-centered cubic (FCC) structure (as shown in Figure 2). SEM analysis of the initial microstructure showed partial recrystallization on the TD, RD, and ND planes (as shown in Figures 3-5). From the TD direction, the non-recrystallized portion appears layered (as shown in Figure 3), with certain dimensional differences between the different layered phases. Furthermore, magnified images of the microstructure reveal fine, sparsely distributed granular second phases. Based on backscatter imaging principles, this can be inferred to be a Mo-rich phase. The volume fraction of the unrecrystallized portion is approximately 27% (Figures 3-5). TEM analysis of the initial microstructure and selected area electron diffraction indicate that the matrix exhibits an FCC structure (Figure 6).
[0072] SEM analysis of the Al alloy specimen subjected to a ~2.2% strain at 750℃ revealed that it retained some recrystallized microstructure characteristics, with no significant decrease in the volume fraction of the non-recrystallized region, and the presence of fine, dispersed granular Mo-rich phases (as shown in Figure 7). TEM analysis of the Al alloy specimen subjected to a ~2.2% strain at 750℃ showed that it exhibited an FCC structure with some weak, short-range ordered diffraction spots. This phenomenon is considered to be the formation of D0 during the ordering process of the Ni-Mo binary alloy. 22 D1 aPrecursor characteristics of the long-range ordered phase. SEM analysis of the tensile fracture specimens of Al alloy at 750℃ revealed significant necking of the fracture surface. Magnified views of the fracture surface showed numerous dimples (Figures 9a and 9b), indicating ductile fracture. Lateral microstructure near the fracture surface revealed that some recrystallized morphology was preserved, the volume fraction of the non-recrystallized region did not decrease significantly, and a certain number of micropores were generated during deformation (Figure 9c).
[0073] Confirmation of the Mo-rich phase structure in the A1 alloy. To clarify the crystal structure of the sparse, dispersed, fine-grained phase (Figures 3-5) observed in the initial microstructure of A1 in the above examples, supplementary verification experiments were conducted. Considering that the size of this phase was too small and the volume fraction insufficient under the original process conditions, making accurate identification by X-ray diffraction and TEM difficult, a sample with the exact same composition as the A1 alloy was prepared, and the same heat treatment temperature was used, but the time was extended to 1 hour to promote the precipitation and coarsening of the Mo-rich phase. SEM and X-ray diffraction tests were performed on this sample. SEM showed that the morphology and compositional contrast of the Mo-rich phase were consistent with those in the A1 alloy (Figure 10). XRD results showed that in addition to the FCC structure, a σ phase structure also existed (Figure 10), indicating that the Mo-rich precipitate phase contained in the partially recrystallized microstructure of the A1 alloy in Example 1 was a σ phase.
[0074] Example 2
[0075] A wide-temperature range (Ni) reinforced with heterostructure 80 Mo 20 ) 96 The preparation method of Al4 alloy differs from that of Example 1 in that the annealing heat treatment temperature and time are 870℃-2min. The resulting alloy is named A2.
[0076] The initial microstructure of the A2 alloy was characterized, and its mechanical properties at room temperature, 500℃, and 750℃ were tested.
[0077] From Example 2 (Ni) 80 Mo 20 ) 96The engineering stress-strain curves of Al4 alloy obtained after tensile tests at different temperatures (as shown in Figure 11) show that its yield strength at -196℃ is approximately 1370 MPa and its total elongation is 28%; at room temperature, its yield strength is approximately 1210 MPa, its tensile strength is approximately 1450 MPa, and its total elongation is approximately 28%; at 500℃, its yield strength is approximately 1060 MPa, its tensile strength is approximately 1340 MPa, and its total elongation is approximately 15%; and at 750℃, its yield strength is approximately 600 MPa, its tensile strength is approximately 660 MPa, and its total elongation is approximately 27%. Under the lower heat treatment temperature and the same treatment time as in Example 1 in Example 2, the volume fraction of recrystallized structure is lower, resulting in an increase in the non-recrystallized region. Therefore, the volume fraction of the non-recrystallized region in Example 2 is greater than 27%. In this embodiment, the room temperature properties of the alloy are significantly better than those in Example 1. Combined with its excellent work hardening ability at room temperature, it is reasonable to infer that its microstructure is predominantly recrystallized. This is because if the microstructure is dominated by non-recrystallized regions, its work hardening ability would significantly decrease, or even disappear entirely. Therefore, the microstructure of this embodiment is predominantly recrystallized, with a non-recrystallized region volume fraction significantly greater than 27%, but less than 50%.
[0078] Comparative Example 1
[0079] A kind of (Ni) 80 Mo 20 ) 96 The preparation method of Al4 alloy differs from that of Example 1 in that the annealing heat treatment temperature and time are 1100℃-5min.
[0080] The alloy prepared in Comparative Example 1 was named A3. The initial microstructure of A3 was characterized, and its mechanical properties were tested at liquid nitrogen (-196℃), room temperature, 500℃, and 750℃. In addition, the microstructure of the tensile fracture specimen at 750℃ was characterized, and the fracture morphology of the tensile fracture surface was observed.
[0081] From Comparative Example 1 (Ni) 80 Mo 20 ) 96The engineering stress-strain curves of Al4 alloy obtained after tensile tests at different temperatures (as shown in Figure 12) reveal the following: at -196℃, the yield strength is approximately 650 MPa, the tensile strength is approximately 1300 MPa, and the total elongation is approximately 63%; at room temperature, the yield strength is approximately 500 MPa, the tensile strength is approximately 1000 MPa, and the total elongation is approximately 64%; at 500℃, the yield strength is approximately 350 MPa, the tensile strength is approximately 900 MPa, and the total elongation is approximately 50%; and at 750℃, the yield strength is approximately 320 MPa, the tensile strength is approximately 470 MPa, and the total elongation is approximately 6%. X-ray diffraction analysis of the initial microstructure shows a face-centered cubic (FCC) structure (as shown in Figure 13). SEM analysis of the initial microstructure reveals fully recrystallized coarse equiaxed crystals with no partial recrystallization characteristics (as shown in Figure 14).
[0082] TEM analysis of the A3 alloy tensile fracture specimen at 750℃ revealed a single-phase FCC matrix structure using selected area electron diffraction (SEM), with no short-range ordered diffraction spots detected (Figure 15). SEM analysis of the A3 alloy tensile fracture specimen at 750℃ showed no obvious necking phenomenon in the fracture morphology (Figure 16a). The fracture morphology was predominantly intergranular fracture, accompanied by a small number of dimples (Figure 16b). Numerous intergranular microcracks were observed on the fracture side, further confirming that the fracture mode was mainly brittle intergranular fracture (Figure 16c).
[0083] Comparative Example 2
[0084] A Ni 80 Mo 20 The preparation method of the alloy differs from that of Example 1 in that: (1) the alloy of Comparative Example 2 has an elemental composition of Ni. 80 Mo 20 (1) No alloying element Al is added; (2) The annealing heat treatment temperature and time are 900℃-4min.
[0085] The alloy prepared in Comparative Example 2 was named A4. The initial microstructure of A4 was characterized, and its mechanical properties were tested at liquid nitrogen (-196℃), room temperature, 500℃, and 750℃.
[0086] From Ni in Comparative Example 2 80 Mo 20The engineering stress-strain curves obtained after tensile tests at different temperatures of the alloy (as shown in Figure 17) reveal the following: at -196℃, the yield strength is approximately 1210 MPa, the tensile strength is approximately 1680 MPa, and the total elongation is approximately 37%; at room temperature, the yield strength is approximately 1020 MPa, the tensile strength is approximately 1300 MPa, and the total elongation is approximately 39%; at 500℃, the yield strength is approximately 900 MPa, the tensile strength is approximately 1120 MPa, and the total elongation is approximately 15%; at 750℃, the yield strength is approximately 580 MPa, the tensile strength is approximately 660 MPa, and the total elongation is approximately 1%. X-ray diffraction analysis of the initial microstructure shows that it has a face-centered cubic (FCC) structure (as shown in Figure 18). SEM analysis of the initial microstructure reveals that it exhibits fine, fully recrystallized equiaxed crystals (as shown in Figure 19).
[0087] Comparative Example 3
[0088] A Ni 80 Mo 20 The preparation method of the alloy differs from that of Example 1 in that: (1) Comparative Example 3 has an elemental composition of Ni. 80 Mo 20 (1) No alloying element Al is added; (2) The annealing heat treatment temperature and time are 1100℃-5min.
[0089] The alloy prepared in Comparative Example 3 was named A5. The initial microstructure of A5 was characterized, and its mechanical properties were tested at liquid nitrogen (-196℃), room temperature, 500℃, and 750℃.
[0090] Ni from Comparative Example 3 80 Mo 20 The engineering stress-strain curves obtained after tensile tests at different temperatures of the alloy (as shown in Figure 20) reveal the following: at -196℃, the yield strength is approximately 600 MPa, the tensile strength is approximately 1360 MPa, and the total elongation is approximately 66%; at room temperature, the yield strength is approximately 460 MPa, the tensile strength is approximately 1010 MPa, and the total elongation is approximately 59%; at 500℃, the yield strength is approximately 340 MPa, the tensile strength is approximately 860 MPa, and the total elongation is approximately 55%; at 750℃, the tensile strength is approximately 460 MPa, and fracture occurs in the elastic stage, exhibiting extreme brittleness. X-ray diffraction analysis of the initial microstructure shows a face-centered cubic (FCC) structure (as shown in Figure 21). SEM analysis of the initial microstructure reveals fully recrystallized coarse equiaxed crystals (as shown in Figure 22).
[0091] Combining the test results of Examples 1-2 and Comparative Examples 1-3 above, it can be seen that the core function of Al microalloying is to suppress harmful phase transformations. For Ni...80 Mo 20 For alloys, regardless of whether the microstructure is coarse or fine-grained, brittle D1 forms at 750℃. a This phase exhibits extreme brittleness, unable to withstand any plastic deformation, as shown in Comparative Examples 2 and 3. In Comparative Example 1, the addition of Al effectively suppressed the harmful long-range order of D1. a The alloy, after phase formation and tensile fracture at 750℃, still exhibits a single-phase FCC state (as shown in Figure 15), which is the fundamental premise for the initial improvement in plasticity (increased to 6%). However, a single compositional design is insufficient to completely solve the brittleness problem. The A3 alloy (Al-doped, equiaxed coarse grains) in Comparative Example 1 still exhibits predominantly intergranular brittle fracture at 750℃, indicating that suppressing harmful phases alone has failed to change its unfavorable deformation mode and fracture mechanism.
[0092] Heterogeneous structure design is a key engineering approach to activate toughness and achieve synergistic strength and plasticity. By introducing partially recrystallized structures on the basis of Al microalloying (Example 1-A1; Example 2-A2), a soft (recrystallized region)-hard (non-recrystallized region) heterostructure was constructed. Taking the Al alloy as an example, only a small number of short-range ordered spots were formed during deformation (Figure 8), indicating that adding Al to the heterostructure suppresses harmful long-range ordered D1. a The phase formation effect remains effective. Furthermore, during deformation, the dispersed σ phase in the Al alloy and the short-range ordered phase formed during deformation synergistically pin each other, ensuring the stability of the heterostructure during high-temperature deformation. Ultimately, a large number of non-recrystallized regions are still observed in the microstructure near the tensile fracture surface, with a volume fraction similar to the initial microstructure (Figure 9c). Thus, at 750℃, the Al alloy achieves high strength (yield strength ~690MPa) while maintaining good plasticity (elongation 29%), and the fracture mode transforms into complete ductile fracture (Figures 9a and 9b). By comparing the mechanical properties of the Al alloy at 750℃ with other alloys reported in the literature, it can be found that it exhibits excellent mechanical properties at 750℃ (as shown in Figure 23). In addition, the Al alloy demonstrates excellent strength-plasticity matching at other temperatures (-196℃, 25℃, 500℃). These results collectively indicate that the heterostructure-strengthened wide-temperature-range Ni-Mo-Al alloy of this invention achieves a good match between strength and plasticity over a wide temperature range.
[0093] The present invention relates to a wide-temperature-range Ni-Mo-Al alloy strengthened by a heterostructure. This heterostructure consists of recrystallized and non-recrystallized regions, achieved through a process of large cold deformation followed by precisely controlled short-time annealing (e.g., 900℃-2min in Example 1 and 870℃-2min in Example 2) on a homogenized solution-solution alloy. The mechanism of heterostructure formation is as follows: cold deformation introduces high-density defects into the alloy, storing sufficient deformation energy. During subsequent annealing, the material first recovers, preferentially triggering recrystallization nucleation in regions with high deformation energy (e.g., original grain boundaries, deformation bands). Due to the short annealing time and optimized temperature, the recrystallization process is not fully completed, thus intentionally retaining a certain proportion of non-recrystallized regions. This results in a heterogeneous microstructure dominated by recrystallized grains and interspersed with banded non-recrystallized regions (observed along the TD and RD directions). The non-recrystallized regions are mainly composed of banded substructures divided by dislocation entanglements and deformation bands, retaining a high dislocation density internally. It should be noted that if the alloy is predominantly composed of unrecrystallized structures, its pre-existing high dislocation density will lead to a significant decrease in work hardening capacity. Furthermore, the precipitation of the σ phase in the alloy also plays a supporting role in regulating recrystallization behavior. During annealing, as the temperature decreases, the solubility of Mo in the Ni-Mo-Al alloy decreases, causing supersaturated Mo to precipitate in an enriched form, forming σ phase particles. Although their volume fraction is low (as in Example 1), these σ phases can still exert a certain pinning effect on grain boundary migration and dislocation reorganization, helping to delay the recrystallization process and maintain the stability of the unrecrystallized region during heat treatment. This makes it easier to construct the desired volume fraction of the unrecrystallized region in the microstructure through different heat treatment processes.
[0094] The method for preparing the wide-temperature-range Ni-Mo-Al alloy of the present invention, by controlling the Al element content and the process parameters of large cold deformation and short-time annealing, achieves the formation of a heterogeneous structure consisting of recrystallized and non-recrystallized regions in the FCC matrix. In the Ni-Mo-Al alloys of Examples 1-2 and Comparative Example 1, the elemental percentage composition is expressed as: (Ni 80 Mo 20 ) 96 Al4, meaning the Al doping content in this embodiment is 4 at.%. As the Al content increases, it further promotes the formation of the σ phase. Therefore, at higher temperatures (such as 1100℃ in Comparative Example 1), the presence of a large amount of σ phase inhibits recrystallization by hindering grain boundary migration, thus preserving the non-recrystallized region and forming the heterostructure of this invention. Therefore, by controlling the Al content and the process parameters of large cold deformation and short-time annealing, the wide-temperature-range Ni-Mo-Al alloy strengthened by the heterostructure in this invention can be obtained.
[0095] The embodiments of the present invention have been described in detail above, but the present invention is not limited to the described embodiments. For those skilled in the art, various changes, modifications, substitutions, and variations made to these embodiments without departing from the principles and spirit of the present invention still fall within the protection scope of the present invention.
Claims
1. A wide-temperature-range Ni-Mo-Al alloy strengthened by a heterostructure, characterized in that, Its elemental percentage composition expression is: (Ni 80+y Mo 20-y ) 100-x Al x Where 2≤x≤8, -5≤y≤2; the Ni-Mo-Al alloy has a face-centered cubic FCC matrix, and the face-centered cubic FCC matrix contains a heterostructure, the heterostructure is composed of recrystallized regions and non-recrystallized regions, wherein the volume fraction of the recrystallized regions is greater than or equal to 50%, and the volume fraction of the non-recrystallized regions is greater than or equal to 20% and less than 50%.
2. The wide-temperature-range Ni-Mo-Al alloy strengthened by heterostructure according to claim 1, characterized in that, The volume fraction of the non-recrystallized region is 20-40%.
3. The wide-temperature-range Ni-Mo-Al alloy strengthened by heterostructure according to claim 1, characterized in that, In the elemental expression of the Ni-Mo-Al alloy, 3≤x≤5, -4≤y≤0.
4. The wide-temperature-range Ni-Mo-Al alloy strengthened by heterostructure according to claim 1, characterized in that, The matrix contains a diffusely distributed σ phase, which is a Mo-rich phase.
5. The wide-temperature-range Ni-Mo-Al alloy strengthened by heterostructure according to claim 1, characterized in that, The non-recrystallized region is layered and mainly consists of strip-shaped substructures divided by dislocation entanglements and deformation bands.
6. A method for preparing a wide-temperature-range Ni-Mo-Al alloy reinforced with a heterostructure as described in claim 1, characterized in that, The process includes the following steps: Step S1, vacuum melting and vacuum casting: Ni, Mo, and Al metal raw materials are vacuum melted and vacuum cast according to the alloy composition ratio to obtain an alloy ingot; Step S2, homogenization treatment; Step S3, cold deformation treatment: multi-pass rolling is performed at room temperature, with a pass rolling amount of 2-6% and a total rolling amount of 60-90%; Step S4, annealing heat treatment: annealing heat treatment is performed in an atmospheric atmosphere to obtain a wide-temperature-range Ni-Mo-Al alloy strengthened by a heterostructure; wherein the annealing heat treatment temperature is 870℃-1150℃ and the holding time is 1.5-20min.
7. The preparation method according to claim 6, characterized in that, Step S1 also includes a surface cleaning pretreatment step for Ni, Mo, and Al metal raw materials before vacuum melting. The surface cleaning pretreatment involves ultrasonic cleaning with anhydrous ethanol, followed by drying. During melting, the Ni, Mo, and Al metal raw materials are stacked according to their melting points, with the order from top to bottom being Mo, Ni, and Al metal raw materials. The melting process is carried out under an argon atmosphere using an electric arc melting method with an arc current of 300-380A. Each melting session lasts for 60-120 seconds. After each melting session, the sample is cooled for 5-10 minutes, then flipped and remelted. The number of remelting sessions is 5-10.
8. The preparation method according to claim 6, characterized in that, In step S2, the homogenization process is carried out in an inert gas atmosphere at 1150℃-1250℃ for 2-24 hours.
9. The preparation method according to claim 6, characterized in that, In step S3, the total rolling volume is 80-90%.
10. The preparation method according to claim 6, characterized in that, In step S2, after homogenization, the material is cooled by water quenching; in step S4, after annealing heat treatment, the material is cooled by water quenching.