A manufacturing method of a nano O-phase reinforced Ti2AlNb alloy with wide temperature range, high strength and high toughness
By controlling particle size and using LDED in-situ manufacturing process, nano-O phase-reinforced Ti2AlNb alloys were prepared, solving the problem of strength and plasticity matching of Ti2AlNb alloys at room temperature and high temperature, and achieving high strength and high toughness over a wide temperature range, which is suitable for aerospace materials.
Patent Information
- Authority / Receiving Office
- CN · China
- Patent Type
- Applications(China)
- Current Assignee / Owner
- GUIZHOU UNIV
- Filing Date
- 2026-04-01
- Publication Date
- 2026-06-02
Smart Images
Figure CN122125239A_ABST
Abstract
Description
Technical Field
[0001] This invention relates to the field of Ti2AlNb alloy manufacturing, and specifically to a method for manufacturing a wide-temperature-range, high-strength, high-toughness nano-O phase-reinforced Ti2AlNb alloy. Background Technology
[0002] Ti2AlNb alloys are valued for their excellent high-temperature performance and moderate density (5-5.7 g / cm³). 3 ), is a nickel-based superalloy used in the aerospace field (8-8.5 g / cm³). 3 Ti2AlNb is a potential alternative material that can effectively reduce structural weight and provide stable service at high temperatures of 600-700℃, showing broad application prospects. Traditional manufacturing processes such as casting, forging, rolling, and powder metallurgy often lead to problems like coarse grains, Nb segregation, and high porosity when preparing Ti2AlNb alloys. Furthermore, these processes are complex and costly, severely hindering the industrial promotion and application of this alloy.
[0003] Additive manufacturing (AM) technology using pre-alloyed powders and mixed powders / wires as raw materials can integrate materials metallurgy and near-net-shape forming, breaking through traditional limitations in cost, cycle time, and composition design, and providing a disruptive solution for forming binary and multi-component difficult-to-machine alloys. Although researchers have prepared Ti2AlNb alloys using additive manufacturing technologies such as L-DED, WAAM, EBAM, and L-PBF, achieving a good balance between strength and plasticity at room temperature and high temperature remains a key bottleneck. Existing research shows that microstructure and mechanical properties exhibit a clear divergence trend. On the one hand, high heat input and moderate cooling rate processes such as L-DED, WAAM, and EBAM tend to form typical Widmanstätten structures, characterized by a large number of micron-sized lath-like O / α2 precipitates distributed in the B2 phase matrix, and continuous brittle α2 phases at grain boundaries. Such microstructures possess high room-temperature strength (e.g., UTS > 1000 MPa), but the continuous α2 phase at grain boundaries easily induces intergranular fracture, severely limiting the material's plasticity (elongation is typically less than 3%). On the other hand, the SLM process, characterized by extremely high cooling rates, can effectively suppress the formation of brittle α2 phases and promote the precipitation of nanoscale O phases. Studies by Zhou et al. and Yang et al. have shown that SLM-prepared alloys are mainly composed of a B2 matrix reinforced by ultrafine nano-O phases (approximately 20 nm). Nano-O phase-reinforced Ti2AlNb alloys exhibit excellent strength and plasticity at room temperature, but their mechanical properties deteriorate significantly at high temperatures (e.g., tensile strength drops to approximately 365 MPa at 650°C, and elongation is only 0.41%). This thermal instability is mainly attributed to the weakening of grain boundary bonding under thermal exposure conditions and the reduction of the pinning effect of unstable nanophases.
[0004] In summary, a key performance trade-off exists in current research: coarse microstructures reduce ductility at room temperature, while nanostructures are insufficient to maintain stability at high temperatures. Therefore, exploring a microstructure control strategy that can both leverage the toughening effect of the nano-O phase and suppress the formation of the grain boundary brittle α2 phase is a crucial problem that urgently needs to be solved. Summary of the Invention
[0005] In view of the shortcomings of the existing technology, the purpose of this invention is to provide a method for manufacturing a wide-temperature-range, high-strength, and high-toughness nano-O phase-reinforced Ti2AlNb alloy.
[0006] To achieve the above objectives, the present invention provides the following technical solution: A method for manufacturing a wide-temperature-range, high-strength, and high-toughness nano-O phase-reinforced Ti2AlNb alloy includes the following steps: S1: Elemental Ti, Al, Nb, and Mo powders prepared by electrode-induced melting gas atomization; S2: Ti powder with a particle size of 53-65μm, Al powder with a particle size of 40-50μm, Nb powder with a particle size of 40-55μm, and Mo powder with a particle size of 15-53μm were screened through a sieve. S3: The selected Ti, Al, Nb and Mo powders are mixed in a planetary ball mill, and then dried in a vacuum drying oven at 120°C for 120 minutes. S4: Use sandpaper to polish the substrate surface to remove the oxide layer, clean with alcohol and acetone and dry to remove residual impurities; S5: Place the dry substrate below the LDED in-situ manufacturing device, and introduce the mixed powder of Ti, Al, Nb and Mo into the LDED in-situ manufacturing device through H2. Deposit the powder layer by layer at a powder feed rate of 5.6 g / min to 70 layers to obtain an O-phase reinforced Ti2AlNb alloy.
[0007] Preferably, in step S1, the Ti, Al, Nb, and Mo powders are screened as spherical elemental powders, wherein the purity of the Ti powder is 99.8%, the purity of the Al powder is 99.9%, the purity of the Nb powder is 99.9%, and the purity of the Mo powder is 99.7%.
[0008] Preferably, in step S2, the average particle size of Ti powder is 60.8 μm, the average particle size of Al powder is 44.7 μm, and the average particle size of Nb powder is 47.5 μm.
[0009] Preferably, in step S3, the ratio of Ti, Al, Nb and Mo powders entering the planetary ball mill is 46:11:43:0.48; the planetary ball mill speed is 100 r / min; and the mixing time is 180 minutes.
[0010] As a preferred option, in step S5, the laser source used in the in-situ manufacturing of LDED has a wavelength of 1064nm, a spot diameter of 3.0mm, a laser power of 1500W, a laser scanning rate of 7mm / s, and an IPG YLS-10000 laser system.
[0011] Preferably, the thickness of each deposition layer in step S5 is 0.4 mm.
[0012] This invention, by adopting the above technical solutions, has significant technical effects: This invention uses Ti, Al, and Nb elemental powders with controlled particle size and trace amounts of Mo as raw materials to in-situ manufacture a nano-O phase-reinforced Ti2AlNb-based alloy. Compared with alloys prepared using pre-alloyed powders, the non-equilibrium solidification of the elemental powders provides favorable nucleation conditions for the precipitation of the nano-O phase. In addition, the negative mixing enthalpy effect and heterogeneous nucleation caused by the addition of Mo result in a unique microstructure: a large amount of nano-sized O phase precipitates on the B2 phase matrix, the brittle α2 phase content is suppressed to 0.2%, and the B2 grain size is refined by 38.4%. The average tensile strength (UTS) at room temperature and high temperature (EL) are 1051.7 MPa and 10.46%, and 673.1 MPa and 9.67%, respectively. The improved plasticity is due to the deformation coordination ability of the nano-O phase and the crack-resistant effect of the discontinuous α2 phase particles at the B2 grain boundaries. This invention provides a key technical path and theoretical support for the low-cost near-net-shape forming of high-performance Ti2AlNb alloys in the aerospace field. Attached Figure Description
[0013] Figure 1 SEM morphology and particle size distribution images of Ti, Al, Nb, and Mo powders, as well as Ti2AlNb pre-alloyed powder, were obtained for this invention.
[0014] Figure 2 This is a comparison diagram of the microstructures of TAN and TAN-0.5Mo samples in this invention.
[0015] Figure 3 This is a schematic diagram of the deposited TAN and TAN-0.5Mo alloy in this invention.
[0016] Figure 4 The image shows the microstructure characterization of the TAN and TAN-0.5Mo samples in this invention.
[0017] Figure 5 This is an elemental distribution diagram of the TAN and TAN-0.5Mo samples in this invention.
[0018] Figure 6 This is a high-resolution EBSD image of the TAN and TAN-0.5Mo samples in this invention.
[0019] Figure 7 This is a comparison chart of the tensile and mechanical properties of TAN and TAN-0.5Mo samples at room temperature and high temperature in this invention.
[0020] Figure 8 The diagram shows the cooling rate test graph, thermal cycle curve, and cooling rate curve of the TAN and TAN-0.5Mo samples in this invention.
[0021] Figure 9 This is a diagram showing the compositional distribution characteristics of Nb-rich microregions at different heights in the TAN and TAN-0.5Mo deposited samples of this invention.
[0022] Figure 10 This is a schematic diagram of the TAN-0.5Mo tissue evolution mechanism in this invention.
[0023] Figure 11 This diagram shows the grain structure of the TAN alloy in this invention, the grain refinement induced by Mo particles, and the grain growth curve induced by enthalpy.
[0024] Figure 12 The images show the room temperature tensile fracture morphology of TAN and TAN-0.5Mo samples in this invention.
[0025] Figure 13 The fracture morphology of TAN and TAN-0.5Mo samples under high temperature tensile conditions of 650℃ in this invention is shown.
[0026] Figure 14 This is an EBSD analysis diagram of the fracture side of the TAN-0.5Mo sample after tensile testing at 25℃ and 650℃ in this invention.
[0027] Figure 15 The images show TEM analysis of TAN and TAN-0.5Mo samples in this invention.
[0028] Figure 16 This is a TEM analysis image of the TAN-0.5Mo sample in this invention.
[0029] Figure 17 This is a phase distribution diagram of the TAN-0.5Mo deposition sample in this invention. Detailed Implementation
[0030] To make the objectives, technical solutions, and advantages of this invention clearer, the invention will be further described in detail below with reference to the accompanying drawings and embodiments. It should be understood that the specific embodiments described herein are merely illustrative and not intended to limit the invention. Furthermore, the technical features involved in the various embodiments of this invention described below can be combined with each other as long as they do not conflict with each other.
[0031] Example 1: Please refer to Figure 1-17 The present invention provides a method for manufacturing a wide-temperature-range, high-strength, and high-toughness nano-O phase-reinforced Ti2AlNb alloy, comprising the following steps: S1: Elemental Ti, Al, Nb, and Mo powders prepared by electrode induction melting gas atomization; the Ti, Al, Nb, and Mo powders are spherical elemental powders, wherein the purity of Ti powder is 99.8%, Al powder is 99.9%, Nb powder is 99.9%, and Mo powder is 99.7%. S2: Ti powder was screened to obtain particles with a diameter of 53-65 μm and an average particle size of 60.8 μm; Al powder had a particle size of 40-50 μm and an average particle size of 44.7 μm; Nb powder had a particle size of 40-55 μm and an average particle size of 47.5 μm; and Mo powder had a particle size of 15-53 μm. The SEM morphology of the screened Ti, Al, Nb, and Mo powders is shown below. Figure 1 As shown in (a)-(d), where (a)Ti, (b)Al, (c)Nb, and (d)Mo have particle size distributions as follows: Figure 1 As shown in (e)-(h), where (e)Ti, (f)Al, (g)Nb, and (h)Mo are mixed powder particles, the EDS results are as follows: Figure 1 (i) Mixed image, (i1)Ti, (i2)Al, (i3)Nb and (i4)Mo, it can be seen that Ti, Al, Nb and Mo powder particles are uniformly distributed and have good sphericity; S3: The selected Ti, Al, Nb and Mo powders are mixed in a planetary ball mill at a ratio of 46:11:43:0.48. The planetary ball mill speed is 100 r / min and the mixing time is 180 minutes. After mixing, the mixture is placed in a vacuum drying oven at 120℃ and dried for 120 minutes. S4: Polish the substrate surface with sandpaper to remove the oxide layer, then clean with alcohol and acetone and dry to remove residual impurities; the substrate used is a TA15 plate with dimensions of 60mm×50mm×5mm; S5: The dried substrate is placed below the LDED in-situ manufacturing device. Ti, Al, and Nb mixed powder is introduced into the LDED in-situ manufacturing device via H2. In the LDED in-situ manufacturing, the laser source has a wavelength of 1064nm and a spot diameter of 3.0mm. The laser system is IPG YLS-10000. Layer-by-layer deposition is carried out at a laser power of 1500W, a scanning rate of 7mm / s, a layer thickness of 0.4mm, and a powder feed rate of 5.6g / min to produce 70 layers. The manufacturing process is carried out under an argon protective atmosphere: two 70-layer Ti-22Al-24.5Nb-0.5Mo alloys are prepared.
[0032] Comparative Example 1: Ti2AlNb alloyed spherical powder prepared by plasma rotating electrode (PREP) was used as the deposition material, and its SEM morphology is shown in the figure. Figure 1 As shown in (j), the particle size distribution is as follows: Figure 1 As shown in (k), Figure 1 EDS results of (l-l2)Ti2AlNb: (l)Ti, (l1)Al, (l2)Nb. It can be seen that the powder has excellent sphericity and high surface smoothness, and the Ti, Al and Nb elements are evenly distributed in a single powder. Steps S4-S5 were performed using the same matrix and LDED in-situ manufacturing apparatus as in Example 1. The manufacturing process was carried out under an argon protective atmosphere: two 70-layer Ti-22Al-25Nb alloys were prepared using Ti2AlNb alloyed powder.
[0033] The following experimental analysis verifies the performance of the Ti2AlNb alloy manufactured in situ using this LDED.
[0034] The Ti-22Al-24.5Nb-0.5Mo alloy prepared in Example 1 is designated as TAN-0.5Mo, and the Ti-22Al-25Nb alloy prepared in Comparative Example 1 is designated as TAN. The performance of the manufactured Ti2AlNb alloy is verified by experimental analysis below.
[0035] I. Microstructure Characterization of Ti2AlNb Alloy Samples Following the L-DED experiment, samples from the middle section of the thin-walled wall were obtained by wire cutting. After grinding, polishing, and etching with Kroll reagent (HF:HNO3:H2O=1:6:10), the microstructure was observed using a German ZEISS GeminiSEM300 scanning electron microscope (SEM), while a Japanese Rigaku Ultima IV microscope was used for further analysis. Phase composition was determined using X-ray diffraction (XRD) at a scanning rate of 1° / min. EBSD samples were prepared using electropolishing and characterized using an EDA-TSL instrument (USA). The test parameters were set as follows: accelerating voltage 15kV, beam current 1nA, and working distance 14mm. Grain size analysis was performed using a 1μm step scan, and O / α2 phase precipitation characteristic analysis was performed using a fine 0.1μm step scan. All data were processed using AZtecCrystal software. For TEM analysis, the sample thickness was first reduced to below 50μm, and then a 3mm diameter disc was thinned using a precision ion polishing system (GATAn695). Finally, the microstructure and phases were analyzed using a JEM-2100F transmission electron microscope (TEM) (Japan).
[0036] like Figure 2 As shown in (a) and (b), the microstructures of TAN and TAN-0.5Mo composite materials are compared: the microstructure of TAN alloy is a typical Widmanstätten structure, with micron-sized needle-like O / α2 phases precipitated in the grains and at the grain boundaries of the B2 matrix; compared with TAN alloy, TAN-0.5Mo alloy not only has finer grains, but also precipitates a large amount of nano-sized O phase, and the content of α2 phase is extremely low.
[0037] II. Mechanical Property Testing of Ti2AlNb Alloy Samples The deposited TAN and TAN-0.5Mo alloy, as shown Figure 3 As shown in (a) and (b), the surface has a metallic luster and no obvious cracks. Nine tensile samples were cut from each of the four thin-walled sections of the TAN and TAN-0.5Mo alloys using wire cutting. These samples were used for tensile property testing at room temperature (25℃) and high temperature (650℃). The samples were sheet-like, with dimensions of 23 mm in length, 7.5 mm in width, and 1.2 mm in thickness. Figure 3 (c) The tensile test was conducted on a Chinese KPL-Focal U100 instrument, in accordance with the Chinese standard GB / T228.1-2010 [Metallic Materials - Tensile Testing at Room Temperature], with a strain rate of 1×10⁻³s⁻¹ and a tensile rate of 0.5 mm / min. The strain was measured using a video extensometer, and the final tensile properties were taken as the average of three samples. The fracture morphology after the tensile test was analyzed by SEM, while the texture and deformation behavior were analyzed by EBSD.
[0038] Figure 4(a) and (b) show the microstructures of TAN and TAN-0.5Mo samples, respectively (the right image is an enlarged view). Figure 4 (c) The XRD analysis results show that in the TAN sample, the phase composition is B2+O+α2, where the B2 phase is the matrix phase and the O phase and α2 phase are intertwined and cannot be directly identified by SEM. Its morphology needs to be further confirmed by EBSD analysis. In the TAN-0.5Mo sample, a large number of nano-sized black particles precipitated on the matrix can be observed under SEM. This structure is different from that of Ti2AlNb alloy prepared by traditional process. XRD analysis shows that the TAN-0.5Mo sample mainly presents the characteristic peak of B2 phase. It is worth noting that the (002)O phase diffraction peak was detected near the 37.5° diffraction angle (as shown in the magnified figure). This phenomenon is consistent with the XRD results of Lin et al. precipitating nano-O phase in Ti2AlNb alloy by hot pressing test.
[0039] To further determine the type and morphological characteristics of the precipitated phase, TEM analysis was performed on the corresponding region of the TAN-0.5Mo sample, revealing nanoscale needle-like precipitates on the matrix. Figure 4 (d) Selected area electron diffraction (SAED) was performed on the matrix, needle-like precipitates, and their interface, respectively. The results are as follows: Figure 4 (e)-(g) SAED pattern diagrams show: (e) B2 phase, (f) O phase, and (g) B2 and O phase interfaces, indicating that the matrix is B2 phase and the nano-sized needle-like precipitates are O phase, with an orientation relationship of
[100] B2 / /
[110] O. Furthermore, TEM energy-dispersive X-ray spectroscopy (EDX) analysis shows ( Figure 4 (h)-(h3)EDX results: (h)Ti, (h1)Al, (h2)Nb and (h3)Mo). Al element is enriched in the O phase, while Nb element is enriched in the B2 phase. This indicates that during the precipitation of O phase from B2 phase, Al element diffuses from B2 phase to O phase. This phenomenon further confirms that the nanoprecipitates are O phase.
[0040] Figure 5 EBSD analysis and elemental distribution results for the TAN and TAN-0.5Mo alloy. Figure 5 As shown in (a) and (b), the grain size of TAN-0.5Mo is significantly smaller than that of TAN; statistical results ( Figure 6(c) shows that the average grain sizes of the two are 230.5 μm and 142.0 μm, respectively. The TAN-0.5Mo grains are refined by 38.4%, and the grain size of TAN-0.5Mo is less than 400 μm, while some grains in TAN exceed 400 μm (greater than 0.4 mm layer thickness), indicating that the grains are formed through epitaxial growth and penetrate multiple melt pool boundaries. Simultaneously, the Kernel Average Misorientation (KAM) of both was analyzed, and the results are as follows: Figure 5 As shown in (d) and (e), the KAM of the TAN-0.5Mo sample (0.61°) is higher than that of the TAN sample (0.54°), and the higher KAM value indicates that the dislocation density (GND) in the sample is higher
[25] . Figure 5 The pole figures and inverse pole figures of (f) and (g) show that TAN has obvious bright spots in the {100}, {110} and {111} pole figures, indicating a strong texture, and a strong (001) preferred orientation in the inverse pole figure; while the crystal plane normal and grain orientation of TAN-0.5Mo are evenly dispersed, and the texture intensity is relatively low, indicating that its preferred orientation is relatively weak.
[0041] To observe the morphology of the precipitates, high-resolution EBSD with a step size of 0.1 μm was used to compare and analyze the microstructure, precipitate texture orientation, and volume fraction of the two alloys. The results were obtained from band-contrast (BC) images. Figure 6 As can be seen from (a) and (b), there are obvious micron-sized needle-like precipitates in the TAN sample, while no obvious precipitates were observed in TAN-0.5Mo. Figure 6 (c)-(d) show the phase composition diagrams of TAN (B2 phase is green, O phase is purple, and α2 phase is yellow). Statistical analysis using AZtecCrystal software shows that the contents of B2 phase, O phase, and α2 phase are 43%, 24.8%, and 31.6%, respectively. Figure 6 (e) It can be seen that a large amount of O / α2 phase precipitates inside the B2 phase grains, while continuous α2 phase mainly precipitates at the grain boundaries. Since the α2 phase has fewer movable slip systems during deformation, this may be the reason for the low plasticity of TAN. For the TAN-0.5Mo sample, a large amount of nanoscale O phase precipitates inside the B2 phase grains, and no obvious α2 phase is observed; although α2 phase precipitates at the grain boundaries, its content is extremely low and discontinuous. Figure 6 (d) The contents of B2 phase, O phase, and α2 phase were 58.3%, 41.5%, and 0.2%, respectively. IPF image ( Figure 6 (f) and (g) show that the precipitates within the grains of both alloys exhibit distinct orientations. Pole figure analysis ( Figure 6(h) and (i) indicate that the three phases B2, O, and α2 in TAN follow the orientation relationship of (001)O / / {0001}α2 and (001)O / / {110}B2, while in TAN-0.5Mo they follow the orientation relationship of (001)O / / {1010}α2 and (001)O / / {110}B2. The O phase and B2 phase both conform to the classical Burgers orientation relationship, indicating that the O phase tends to nucleate along the grain boundary or inside the B2 phase grain with a specific orientation relationship. This phenomenon is consistent with the research conclusions of Zhou et al.
[0042] Tensile test results show that, at room temperature (25°C), the average UTS and average EL of TAN and TAN-0.5Mo alloys are 895.6 MPa and 1.23%, and 1051.7 MPa and 10.46%, respectively (see...). Figure 7 (a) and (b)). The TAN alloy exhibits brittle fracture. Compared to the TAN alloy, the TAN-0.5Mo alloy shows a 17% increase in room temperature UTS. Notably, its elongation at high temperatures (EL) is significantly improved, and the fracture mode shifts to ductile fracture. At high temperatures (650°C), the average UTS and EL of both are 677.4 MPa and 0.76%, and 673.6 MPa and 9.67%, respectively, demonstrating a significant improvement in high-temperature elongation for TAN-0.5Mo.
[0043] The manufacturing processes for Ti2AlNb alloys are diverse. To clarify the position of the multi-principal element in-situ manufacturing process among commonly used processes, its mechanical properties are compared with those of components prepared using other processes. The results are as follows: Figure 7 As shown in (c) and (d). Figure 7 (c) Comparison of mechanical properties of common processes at room temperature (25°C): Most traditional manufacturing processes (HPS, HIP, PM, SPS) are difficult to achieve a good match between strength and plasticity, while TAN-0.5Mo's UTS and EL surpass most traditional processes, achieving an excellent match between strength and plasticity at room temperature; compared with additive manufacturing processes (SLM, WAAM, EBAM), TAN-0.5Mo's mechanical properties are also better, especially in terms of elongation. Conventional additive manufacturing processes generally have poor plasticity, and although the EL of the SLM process can be improved through post-processing, the UTS will decrease accordingly. Figure 7(d) shows the comparative results of the mechanical properties of Ti2AlNb alloys at 650℃. The UTS and EL of TAN-0.5Mo are significantly better than those of alloys prepared by traditional processes. Although the Ti2AlNb alloys prepared by SLM technology exhibit better room temperature strength and ductility, their strength and ductility decrease sharply under high-temperature tensile stress. Even after heat treatment, their high-temperature ductility is significantly improved, but TAN-0.5Mo prepared in this study still has the advantage in UTS and EL. In summary, Ti2AlNb alloys manufactured in situ using multi-principal element methods have high engineering application value. Although the mechanical properties of this alloy at room temperature and high temperature are still inferior to those of heat-treated forgings, its room / high-temperature strength and ductility meet the service requirements. Furthermore, the forging process combined with subsequent heat treatment suffers from poor flexibility and high cost, which limits its application in the manufacture of Ti2AlNb alloys.
[0044] III. Microstructure Evolution Analysis of Ti2AlNb Alloy Samples Use an infrared camera to obtain TAN ( Figure 8 (a)) and TAN-0.5Mo( Figure 8 (b) Temperature curves at the midpoints (points A1 and A2) of the first layer and the midpoints (points B1 and B2) of the 35th layer of the sample, as shown in the figure. Figure 8 As shown in (c), combined with the pseudo-binary phase diagram of the Ti-22Al-xNb alloy ( Figure 8 (f) The analysis is as follows: The thermal cycling temperature of the first 15 layers of the TAN sample is above 1060℃. In this temperature range, the O phase and α2 phase will not precipitate. The thermal cycling temperature of the next 3 layers crosses 1060℃ (the lowest temperature is below this value and the highest temperature is above this value). The precipitation and re-dissolution of α2 phase alternate. However, because the re-dissolution temperature is higher than the precipitation temperature, the precipitation of α2 phase is not obvious. The thermal cycling temperature of the next 6 layers is between 880-1060℃. The O phase and α2 phase begin to precipitate stably. The thermal cycling temperature of the last 10 layers is below 880℃. Only the O phase continues to precipitate. Finally, the TAN alloy forms a phase composition of B2+O+α2. The enthalpy of mixing between Ti, Al, Nb and Mo is shown in Table 1.
[0045] Elements Ti Al Nb Mo Ti - -30 2 -4 Al - -18 -5 Nb - -6 Mo - The results indicate that Ti-Al exhibits a significant negative enthalpy of mixing (-30 kJ / mol), while Ti-Nb exhibits a positive enthalpy of mixing (2 kJ / mol). This thermodynamic difference drives Ti and Al to preferentially combine to form stable Ti-Al clusters, thereby inducing the formation of local Nb-rich microregions and creating favorable nucleation conditions for O phase precipitation. Simultaneously, the high cooling rate of L-DED retains the Nb-rich microregions in the liquid phase within the B2 phase matrix. Correlation analysis using SEM and EDS (…) Figure 9Compositional distribution characteristics of Nb-rich microregions at different heights in the deposited sample: (ad) Top region: (a) low-magnification SEM microstructure, (b) magnified view of the corresponding area (a) with yellow circle, (c) surface scan results and (d) line scan results of Line 1; (eh) Middle region: (e) low-magnification SEM microstructure, (f) magnified view of the corresponding area (d) with yellow circle, (g) surface scan results and (h) line scan results of Line 2) revealing the formation mechanism of the nano-O phase. Figure 9 (a) shows a backscattered electron (BSE) image of the top region of the deposited sample, revealing a distinct Nb-rich microregion with no obvious precipitates inside, mainly composed of the B2 phase. Figure 9 (b)). EDS surface scan ( Figure 9 (c) and line scan ( Figure 9 (d) confirms that the Nb concentration in this microregion is high. This is because the top region experienced fewer in-situ thermal cycles, resulting in insufficient elemental diffusion motive force. The high Nb concentration acts as a strong β stabilizer, thereby inhibiting the decomposition of the B2 phase. In contrast, although Nb-rich microregions also exist in the middle of the deposited sample, the internal Nb concentration is relatively low due to sufficient in-situ thermal cycling that drove elemental diffusion. Figure 9 (g) and (h)). It is worth noting that the tissue characteristics inside and outside the Nb-rich microregions were compared ( Figure 9 (e) and (f) revealed that the O phase generated inside the microregion exhibited a high-density nanoscale morphology, with a size significantly smaller than the precipitated phase outside the microregion. This feature confirms the dominant role of the solute dragging mechanism: the high concentration of Nb solute at the center of the Nb-rich microregion increased the phase boundary migration resistance, effectively inhibiting grain growth while promoting a high nucleation rate, thus forming a unique nanoscale O phase structure.
[0046] A schematic diagram of the TAN-0.5Mo tissue evolution mechanism is shown below. Figure 10 As shown, the premixed powder melts rapidly under the action of the laser, and the temperature rises rapidly to its peak value; due to the non-equilibrium solidification caused by the mixing enthalpy, Ti and Al preferentially form Ti-Al intermetallic compounds, which in turn induces the formation of Nb-rich microregions. Figure 10 (b) Meanwhile, due to the rapid cooling effect of additive manufacturing, the B2 phase does not have time to precipitate the α2 and O phases, resulting in the Ti2AlNb alloy phase composition that has not undergone thermal cycling being a single B2 phase, and the Nb-rich microregions are preserved. Figure 10 (c)). By Figure 9 (c) It can be seen that the temperature of the first 9 layers in the middle part (point B2) of the TAN-0.5Mo sample spans 1060℃, during which a small amount of α2 phase precipitates along the grain boundaries; in the subsequent 6 layers (temperature 880-1060℃), O phase and α2 phase begin to precipitate, with α2 phase nucleating along the grain boundaries, and a large amount of extremely fine nano-O phase precipitating inside the Nb-rich microregions. Figure 10 (d)); The subsequent 22-layer thermal cycling temperature was below 880℃ (the number of layers in this range was significantly greater than that of the TAN sample), which promoted the precipitation of a large amount of O phase. Figure 10 (e) Furthermore, trace amounts of Mo can effectively refine the size of the O / α2 precipitates and the B2 matrix phase in the Ti2AlNb alloy, while reducing the continuity of the α2 phase. The cooling rates of TAN and TAN-0.5Mo were calculated using thermal cycling curves, and the results are as follows: Figure 9 In (d) and (e), the cooling rates of the first layer of TAN and TAN-0.5Mo samples were 2144℃ / s and 1506℃ / s, respectively; the cooling rates of the 35th layer were 1425℃ / s and 1068℃ / s, respectively. It can be observed that the cooling rate of the TAN-0.5Mo sample was consistently higher than that of TAN, both at the midpoint of the first and 35th layers. The faster cooling rate, combined with the synergistic effect of trace amounts of Mo, further refined the size of the O / α2 precipitates and reduced the continuity of the α2 phase, ultimately resulting in a phase composition dominated by B2 and nano-O phases in the TAN-0.5Mo sample, with only a small amount of discontinuous α2 phase precipitated at the grain boundaries.
[0047] The B2 grain refinement mechanism of the TAN-0.5Mo sample mainly stems from the rapid cooling rate, the effect of trace amounts of Mo, and the mixing enthalpy effect. EBSD analysis (…) Figure 3 (c) It can be seen that, compared with the TAN sample, the B2 grain refinement of TAN-0.5Mo reaches 38.4%; among them, some grains in the TAN sample have a size greater than 400 μm (layer thickness of 0.4 mm), indicating the presence of some columnar crystals. Figure 11 (a)), while the grain size of the TAN-0.5Mo samples was all less than 400 μm, and they were basically equiaxed grains. In the additive manufacturing process, the temperature gradient G and the solidification rate V jointly determine the grain growth along the forming direction, while the cooling rate (G×V) is the core factor affecting the microstructure size. Previous studies have shown that as the cooling rate increases, the microstructure size decreases accordingly. Figure 8 The analysis results in (d) and (e) show that the TAN-0.5Mo sample has a faster cooling rate, resulting in a smaller B2 grain size. From the perspective of element melting and nucleation mechanisms, when multiple principal elements enter the molten pool, Ti, Al, and Nb melt first and form a Ti2AlNb alloy because their melting points are lower than Mo's. Mo, which melts slightly later, enters the molten pool and acts as a solid-phase particle for heterogeneous nucleation, thus promoting heterogeneous nucleation within the molten pool. According to the heterogeneous nucleation theory, the formula for calculating the nucleation work is:
[0048]
[0049] In the formula, ΔGk For the critical nucleation work, r k Let σ be the critical nucleation radius, σ be the surface energy between the heterogeneous nucleation core (crystal nucleus) and the liquid phase, and θ be the energy between the crystal nucleus and the solid phase particles. m For latent heat of fusion, T m Let ΔT be the theoretical crystallization temperature, and ΔT be the difference between the theoretical and actual crystallization temperatures. The critical nucleation work ΔG for heterogeneous nucleation is... k The critical nucleation work is mainly determined by the wetting angle θ: the smaller θ is, the lower the critical nucleation work. The size of the wetting angle θ depends on the difference in crystal structure between the nucleated new phase and the solid phase particles. Since the added Mo has the same body-centered cubic structure as B2, and its lattice parameters are very close to those of the B2 phase, the difference in crystal structure between the nucleus and the solid phase particles is minimal, allowing for near-complete wetting. In this case, the critical nucleation work required for nucleation is significantly reduced, thereby effectively promoting heterogeneous nucleation within the melt pool, ultimately driving the transformation of columnar crystals to equiaxed crystals and achieving grain refinement of B2. Figure 11 (b)).
[0050] In laser in-situ manufacturing, the premixed Ti, Al, Nb, and Mo powders undergo melting and metallurgical reactions. The accompanying thermal effects of these two processes alter the heat flow direction, thus affecting the epitaxial growth of grains. Therefore, it is necessary to calculate the enthalpy of mixing of the liquid alloy of the Ti, Al, Nb, and Mo system. Here, the quaternary alloy extrapolation method proposed by Ouyang et al. is used to calculate the enthalpy of mixing of this system. The calculation results show that the enthalpy of mixing of the Ti, Al, Nb, and Mo premixed powders is -28.25 kJ / mol. This enthalpy of mixing is less than zero, indicating that the alloying process is an exothermic reaction. The heat released by the reaction changes the heat flow direction in the local region at the solid / liquid interface front. When the next layer is deposited, the epitaxial growth of the original grains along the deposition direction is interrupted, thereby promoting the nucleation and growth of new grains at the solid / liquid interface front, ultimately achieving microstructure refinement (e.g., ...). Figure 12 (c) is shown.
[0051] The tensile fracture surfaces of TAN and TAN-0.5Mo samples at room temperature are as follows: Figure 12 As shown (where (a,b) are overall fracture images: (a) TAN and (b) TAN-0.5Mo; (c,d) are fracture characteristic regions of TAN: (c) intergranular fracture and (d) transgranular fracture; (e,f) are fracture characteristic regions of TAN-0.5Mo: (e) ductile fracture and (f) dimples). The TAN sample exhibits typical brittle fracture characteristics under room temperature tensile conditions, with intergranular and transgranular fracture modes being the predominant modes. Figure 12 (a) In the intergranular fracture region, a rock-patterned structure formed by coarse equiaxed grains can be clearly observed, and obvious cracks exist between the grains. Figure 12 (c) indicates that cracks tend to propagate along the B2 grain boundaries during tensile testing; in the transgranular fracture region ( Figure 12 (d) shows numerous typical river patterns, with their orientation perpendicular to the grain boundaries. For the TAN-0.5Mo alloy, obvious necking can be observed on the room-temperature tensile fracture surface. Figure 12 (b) Magnified observation of the local area reveals cracks and dimples on the fracture surface. Figure 12 (e)-(f)) indicates that the fracture mode is ductile fracture. From the correlation between microstructure and mechanical properties: the average grain size of the B2 phase in the TAn alloy reaches 230.5 μm. The coarse grain structure reduces the number of grain boundaries that can hinder crack propagation, resulting in low elongation and brittle fracture characteristics. Furthermore, a large amount of α2 phase precipitation can be observed in the TAn sample (see...). Figure 6 (c) The α2 phase has fewer activatable slip systems during room temperature stretching and tends to precipitate continuously at grain boundaries, further significantly reducing the alloy elongation. The B2 grain size of the TAN-0.5Mo alloy is effectively refined, and the increased grain boundaries can significantly block crack propagation. At the same time, the precipitation of brittle α2 phase is suppressed, with only a small amount and discontinuous distribution at grain boundaries. A large amount of nanoscale O phase precipitates in the B2 phase. Previous studies have shown that compared with the α2 phase, the O phase has a higher bonding strength with the B2 phase at room temperature. The nanoscale O phase can achieve synergistic enhancement of alloy strength and ductility, which is the key reason why the TAN-0.5Mo alloy exhibits ductile fracture.
[0052] Figure 13 The fracture morphology of TAn and TAN-0.5Mo samples under high-temperature tensile conditions at 650℃ is shown. The high-temperature tensile fracture of the TAn sample mainly exhibits intergranular fracture, which is a typical brittle fracture. Figure 13 (a)); Fracture analysis shows that the crack still tends to propagate along the B2 phase grain boundaries ( Figure 13 (c) indicates that the coarse B2 phase grains remain the key factor leading to the low high-temperature tensile elongation of the TAN sample. From a microscopic perspective, the lack of cross-slip ability of the α2 phase easily leads to dislocation accumulation at grain boundaries and stress concentration, making the grain boundaries where the α2 phase is located crack initiation sites; moreover, the continuous precipitation of the α2 phase along the grain boundaries further severely degrades the high-temperature tensile properties of the Ti2AlNb alloy, which is consistent with the brittle fracture characteristics of the TAN sample. For the TAN-0.5Mo sample, no obvious cracks were found on the high-temperature tensile fracture surface at 650℃. Figure 13 (b)) The fracture morphology simultaneously exhibits cleavage steps and dimples (ductile characteristics) (e.g. Figure 13 As shown in (d) and (e), the fracture mode is a mixed fracture mode. It should be noted that during the high-temperature tensile process at 650℃, some of the B2 phase will transform into the O phase. Since the number of movable slip systems in the O phase is less than that in the B2 phase, its high-temperature tensile plasticity is slightly lower than that at 25℃.
[0053] Figure 14 The EBSD analysis results of the fracture surface of TAN-0.5Mo alloy after tensile testing at 25℃ and 650℃ are presented, along with IPF plots ( Figure 14 As shown in (a) and (e), during tensile testing at 25℃, the grain orientation of alloy B2 is concentrated, mainly distributed along the direction parallel to the tensile direction (X direction). The grains elongate along the X direction, and the color of individual grains is uneven, indicating that the grain orientation changes during deformation. After tensile testing at 650℃, the grain orientation distribution becomes more uniform, and no obvious elongation phenomenon is observed. The analysis of the grain boundary characteristic diagram shows that ( Figure 14 (b) and (f)) After stretching at 25℃, the proportion of high-angle grain boundaries (HAGBs) in the alloy reached as high as 89.7%, while the proportion of low-angle grain boundaries (LAGBs) was only 10.3%, with no obvious traces of grain boundary migration. After stretching at 650℃, the proportion of LABs significantly increased to 42.6%, while the proportion of HAGBs decreased to 57.4%, with obvious traces of migration and subgrain boundary formation appearing at the grain boundaries. This phenomenon is attributed to the enhanced atomic diffusion ability of grain boundaries at high temperatures, the activation of dynamic recovery and recrystallization processes, and the participation of grain boundaries in deformation through slip and subgrain rotation, thereby reducing the hindering effect on deformation. After stretching at 25℃, the KAM value of the alloy was 1.37° ( Figure 14 (c)), significantly higher than the 0.61° of the undeformed state ( Figure 5 (e) A higher KAM value is associated with a higher geometrically necessary dislocation density, and an increase in geometrically necessary dislocation density is beneficial to improving the deformation behavior of the alloy, which is also the reason for the excellent plasticity of the alloy at this temperature. After stretching at 650℃, the KAM value (0.85°) increases slightly. Figure 14 (g) Since the stretching temperature of 650℃ falls within the phase transition range of B2 to O phase, some B2 phase will transform into O phase during the stretching process, resulting in an increase in the proportion of O phase under high temperature conditions. Figure 14 (d)-(i)); To clarify the two-phase texture characteristics, select Figure 14 Analyze the local regions of (a) and (e). Figure 14 (j), (k)), the results show that the orientation relationship between the B2 phase and the O phase is {110}B2 / / {001}O, and the texture strength of the stretched samples at both temperatures is significantly higher than that of the unstretched samples. Figure 14 (g) The formation of this texture is the result of the combined effects of plastic deformation, grain rotation, and temperature. On the one hand, during the stretching process, the crystal undergoes plastic deformation under heating and tensile loads, and the internal atoms rearrange to adapt to the applied stress. On the other hand, grain rotation and deformation further promote the formation of preferred orientations. The main slip system of the B2 phase in the Ti2AlNb alloy is {110}. <111> {112} <111> and {123} <111>
[49] , Schmidt factor analysis showed that ( Figure 14(l), (m)), the mean Schmid factors of the three slip systems of the B2 phase after stretching at 25℃ and 650℃ are 0.48, 0.45, 0.45 and 0.48, 0.45, 0.47 respectively, indicating that {110} <111> As the main slip direction, in addition to the slip system being easily activated at high temperatures, {123} <111> The slip system is also easier to start because {123} <111> The critical analytical shear stress is relatively high, making it difficult to activate at room temperature, but it can be activated at high temperatures, similar to {110}. <111> This forms a double-slip system, ultimately improving the deformation uniformity of the alloy.
[0054] The room temperature and high temperature tensile properties of TAN and TAN-0.5Mo deposited samples were compared. The results showed that, under RT conditions, the average yield strength (YS) of TAN-0.5Mo was increased by 145.1 MPa compared to the TAN sample. Notably, the plasticity of TAN-0.5Mo was significantly improved under both room temperature and high temperature conditions. Therefore, it is necessary to further explore the mechanism of the improved room temperature mechanical properties of TAN-0.5Mo, and the potential mechanism of its improved plasticity under both room temperature and high temperature conditions.
[0055] Microstructural analysis revealed that the B2 grain size of the TAN-0.5Mo alloy was significantly finer and the dislocation density was also increased compared to the TAN sample. These observations indicate that the grain structure must be taken into account. ), (including inherent lattice friction and grain refinement) and dislocation mechanisms ( The strengthening effect produced by solid solution strengthening. Meanwhile, given that Mo acts as an effective B2 phase stabilizer in Ti2AlNb alloys, the effect of solid solution strengthening must also be quantitatively evaluated. Furthermore, the TAN-0.5Mo alloy contains high-density nano-O phase precipitates. These precipitates have been shown to produce a considerable precipitation strengthening effect. This strengthening effect is achieved, in particular, through their interaction with dislocation movement within the B2 matrix. Since dislocations bypass these precipitates rather than shear them, the Orovan bypass mechanism can be used to effectively assess the strengthening effect of these features. By integrating these individual contributions, the theoretical yield strength of the alloy can be predicted and verified against experimental values.
[0056] (1) Refined grain strengthening ( ) The rapid cooling rate, heterogeneous nucleation effect, and mixing enthalpy effect collectively suppressed grain growth in the TAN-0.5Mo sample. Compared to the TAN sample, the grain size of the TAN-0.5Mo sample was refined by 38.4%. Grain refinement leads to the formation of numerous new grain boundaries, which in turn hinder dislocation slip between different grains, ultimately resulting in improved material strength. The difference in strength contribution due to grain boundary strengthening can be calculated using the Hall-Petch formula.
[0057]
[0058] in, This represents the lattice friction stress (274 MPa). k HP It is the Hall-Petch constant, with a value of 313 MPa·um. 0.5 , d m The average grain size (d) of the TAN-0.5Mo deposited matrix. m =142.0μm), and the strength contribution from grain refinement is calculated to be 300.2MPa.
[0059] (2) Orowan reinforcement ( ) To investigate the strengthening mechanism of the nano-O phase, TEM analysis was performed on the interface region between the nano-O phase and the B2 phase. Figure 15 (a) TEM bright-field image. For example... Figure 13 The HRTEM image of region (b) marked in (a) is as follows: Figure 15 As shown in (b), the orientation relationship between the B2 and O phases satisfies the crystallographic relation (110)B2 / / (002)O. Further IFF and geometric phase analysis (GPA) were performed on this region, and the results are as follows: Figure 15 As shown in (c) and (d), a significant strain concentration phenomenon, along with dislocations and lattice distortion, is observed at the B2 / O phase interface. This phenomenon may originate from the effect of Mo, an element that acts as a stabilizing element for the B2 phase and induces lattice distortion at the B2 / O phase interface in the Ti2AlNb alloy. Figure 15 (f) It is evident that dislocations readily coalesce at grain boundaries, forming stacking faults there; these dislocations do not penetrate the nano-O phase but instead become entangled around the periphery of the nano-O phase. This process leads to a large accumulation of dislocations (such as...). Figure 15 (e) shows that the dislocation entanglement phenomenon is eventually formed, which plays an important role in improving the strength of lath alloys containing long O phase.
[0060] In the Orowan strengthening mechanism, the nano-O phase, which is difficult for dislocations to penetrate within the matrix, hinders dislocation movement through its interaction with dislocations, thereby suppressing plastic deformation and ultimately enhancing the strength of the B2 matrix. The contribution of Orowan strengthening to the material strength can be estimated using a formula.
[0061]
[0062] In the formula, G m It is the shear modulus of the B2 phase matrix (37.5 GPa). b It is the Burgers vector (0.2797 nm). d o The average particle size of the nano-O phase is 255.83 nm. Vo The volume fraction of the enhanced nano-O phase was calculated to contribute 509 MPa to the strength due to Orowan reinforcement.
[0063] (3) Solid solution strengthening ) Solid solution strengthening is calculated using the formula: Among them, B i The strengthening coefficient of solute Mo is 575 MPa·at. -2 / 3 ), c i This represents the concentration of solute Mo (0.25 at%). Calculations show that the strength contribution from solid solution strengthening is 10.6 MPa.
[0064] (4) Dislocation reinforcement ) The higher the dislocation density, the more easily they intersect and form cutting steps during their movement, leading to dislocation entanglement. Dislocation entanglement hinders further dislocation slip, ultimately increasing the material's strength. The increment of dislocation strengthening is calculated using the formula:
[0065] in, M This is the average orientation factor. Ti2AlNb is a BCC structure alloy, so its value can be taken as 3.06. β It is the dislocation barrier coefficient (0.35). The geometric dislocation density of TAN-0.5Mo was statistically analyzed using AZtecCrystal software. = 0.76×10 14 / m 2 Calculations show that the strength contribution from dislocation strengthening is 97.9 MPa.
[0066] In summary, Orowan strengthening contributes 49.3% to the overall strength, indicating that the precipitated nano-O significantly improves the mechanical properties of the Ti2AlNb alloy. The theoretical yield strength, derived from the sum of grain boundary strengthening, dislocation density, solid solution, and the Orowan mechanism, is 917.7 MPa. Compared to the experimentally measured yield strength (1032.4 MPa), this theoretical prediction deviates by 11.1%. This reasonable consistency validates the reliability of the theoretical model and confirms that the excellent strength of the TAN-0.5Mo alloy is determined by the synergistic effect of these mechanisms.
[0067] VII. TAN-0.5Mo Plasticity Enhancement The abundant precipitation of brittle α2 phase at the B2 phase grain boundaries in Ti2AlNb alloys is the main reason for the sharp deterioration of room-temperature plasticity. Furthermore, the size effect of the O phase also significantly influences alloy plasticity: nanoscale O phases exhibit excellent deformation compatibility, while coarse micron-sized O phases significantly reduce ductility. Therefore, the improvement in room-temperature plasticity of TAN-0.5Mo alloys is mainly attributed to the effective suppression of α2 phase content and the dispersed precipitation of nanoscale O phases. Figure 16 TEM analysis of the microstructure of TAN-0.5Mo alloy after tensile testing at 650 °C is presented. It can be seen that significant planar slip and cross-slip are activated within the soft matrix B2 phase. Figure 16 (a) indicates that the activation of the multislip system effectively coordinates macroscopic plastic deformation; although dislocations pile up at the O / B2 interface ( Figure 16 (b)), but stress concentration successfully drove dislocations to cut through part of the nano-O phase ( Figure 16 (c) allows stress to be released. It is noteworthy that, although the O-phase slip system is relatively small, it induces a deformation twinning mechanism under high temperature and pressure conditions to further coordinate deformation. Figure 16 (d) High-magnification observation revealed a clear parallel lamellar structure within the twins, exhibiting typical characteristics of nanotwins. Figure 16 (e)). High-resolution TEM characterization of nanotwins ( Figure 16 (f)), combined with IFFT image ( Figure 16 (f1) and εxx strain distribution diagram ( Figure 16 (f2) indicates that numerous dislocations and lattice distortions exist within the twin. This atomic-scale strain concentration and coordination mechanism effectively dissipates deformation energy, preventing stress singularities from forming within the hard phase. Furthermore, from Figure 17 The phase distribution of the deposited sample shows that granular α2 phase precipitated at the B2 grain boundaries of the TAN-0.5Mo alloy. Figure 17(d) Related studies have confirmed that the precipitation of fine-grained α2 phase at the B2 grain boundaries can endow the alloy with better crack resistance, thereby effectively improving the high-temperature plasticity of the alloy.
[0068] This invention utilizes elemental powders of Ti, Al, Nb, and Mo as raw materials to fabricate a high-strength, high-ductility Ti2AlNb alloy in situ using L-DED (Liquid-Distilled Electrode Emulsion). The thermal cycling process, microstructure, and mechanical properties of the deposited samples were systematically analyzed. This work provides new insights into the preparation of low-cost, high-performance Ti2AlNb-based alloys. The main conclusions are as follows: (1) During the in-situ manufacturing process, the non-equilibrium solidification of elemental powders leads to the formation of local Nb-rich microregions, which provides favorable nucleation conditions for the precipitation of nano-O phase. At the same time, the fast cooling rate, the non-uniform nucleation of trace Mo, and the negative mixing enthalpy refine the microstructure of TAN-0.5Mo, forming a Ti2AlNb alloy dominated by B2 phase + nano-O phase, which effectively suppresses the large-scale precipitation of brittle α2 phase.
[0069] (2) The higher room-temperature YS of the TAN-0.5Mo alloy is attributed to a synergistic strengthening effect of multiple mechanisms, with orovan strengthening playing a dominant role (509 MPa), while grain boundary strengthening (300.2 MPa), dislocation strengthening (97.9 MPa), and solid solution strengthening (10.6 MPa) also play a role. The difference between the theoretically predicted total strength and the experimental measurement is only 11.1%, which verifies the accuracy of the proposed strengthening model.
[0070] (3) The improvement in the ductility of the alloy is mainly attributed to the excellent deformation coordination of the nano-O phase and the crack-suppressing effect of the discontinuous and fine α2 phase at the B2 grain boundary. These factors work together to change the fracture mode of the alloy from the traditional brittle fracture to ductile fracture at room temperature, and to a mixed fracture mode at high temperature.
[0071] (4) The prepared TAN-0.5Mo alloy achieves a breakthrough in synergistic strength and plasticity at room temperature and high temperature: the average UTS and average EL at room temperature and high temperature are 1051.7MPa and 10.46% and 673.1MPa and 9.67%, respectively. The overall performance is significantly better than that of traditional processes (HIP, PM, SPS) and mainstream additive manufacturing technologies (SLM, WAAM, EBAM).
[0072] The above description is only a specific embodiment of the present invention, but the technical features of the present invention are not limited thereto. Any changes or modifications made by those skilled in the art within the scope of the present invention are covered by the patent scope of the present invention.
Claims
1. A method for manufacturing a wide-temperature-range, high-strength, high-toughness nano-O phase-reinforced Ti2AlNb alloy, characterized in that, Includes the following steps: S1: Elemental Ti, Al, Nb, and Mo powders prepared by electrode-induced melting gas atomization; S2: Ti powder with a particle size of 53-65μm, Al powder with a particle size of 40-50μm, Nb powder with a particle size of 40-55μm, and Mo powder with a particle size of 15-53μm were screened through a sieve. S3: The selected Ti, Al, Nb and Mo powders are mixed in a planetary ball mill, and then dried in a vacuum drying oven at 120°C for 120 minutes. S4: Use sandpaper to polish the substrate surface to remove the oxide layer, clean with alcohol and acetone and dry to remove residual impurities; S5: Place the dry substrate below the LDED in-situ manufacturing device, and introduce the mixed powder of Ti, Al, Nb and Mo into the LDED in-situ manufacturing device through H2. Deposit the powder layer by layer at a powder feed rate of 5.6 g / min to 70 layers to obtain an O-phase reinforced Ti2AlNb alloy.
2. The method for manufacturing a wide-temperature-range, high-strength, high-toughness nano-O phase reinforced Ti2AlNb alloy according to claim 1, characterized in that, In step S1, Ti, Al, Nb, and Mo powders are screened to be spherical elemental powders, wherein the purity of Ti powder is 99.8%, the purity of Al powder is 99.9%, the purity of Nb powder is 99.9%, and the purity of Mo powder is 99.7%.
3. The method for manufacturing a wide-temperature-range, high-strength, high-toughness nano-O phase reinforced Ti2AlNb alloy according to claim 1, characterized in that, In step S2, the average particle size of Ti powder is 60.8 μm, the average particle size of Al powder is 44.7 μm, and the average particle size of Nb powder is 47.5 μm.
4. The method for manufacturing a wide-temperature-range, high-strength, high-toughness nano-O phase reinforced Ti2AlNb alloy according to claim 1, characterized in that, In step S3, the ratio of Ti, Al, Nb and Mo powders entering the planetary ball mill is 46:11:43:0.48; the planetary ball mill speed is 100 r / min, and the mixing time is 180 minutes.
5. The method for manufacturing a wide-temperature-range, high-strength, high-toughness nano-O phase-reinforced Ti2AlNb alloy according to claim 1, characterized in that, In step S5, the laser source used in the in-situ LDED manufacturing process has a wavelength of 1064nm, a spot diameter of 3.0mm, a laser power of 1500W, a laser scanning rate of 7mm / s, and an IPG YLS-10000 laser system.
6. The method for manufacturing a wide-temperature-range, high-strength, high-toughness nano-O phase-reinforced Ti2AlNb alloy according to claim 5, characterized in that, In step S5, each deposition layer has a thickness of 0.4 mm.