Method for in-situ manufacturing of high-performance Ti2AlNb alloy through particle size control and laser directional energy deposition

Abstract

The invention relates to the field of Ti2AlNb alloy manufacturing, and discloses a method for in-situ manufacturing of a high-performance Ti2AlNb alloy through particle size regulation and laser directional energy deposition, particle size regulation and premixing of Ti powder, Al powder and Nb powder are carried out through a planetary ball mill, and the Ti2AlNb alloy with excellent performance is manufactured in situ through LED. According to the method, the three kinds of particles can be fully melted and mixed, so that element enrichment and depletion are reduced, meanwhile, phase compositions of the top area and the middle area of a component are B2 and B2 + alpha2 + O phases respectively, and the cooling rate is increased through uneven heat conduction; the Nb particles with the high melting point serve as heterogeneous nucleation substrates, the local heat flow direction in front of a solid / liquid interface is changed through heat released by negative mixing enthalpy, and therefore the growth direction of original crystal grains is changed, and finally the structure is refined. The average tensile strength at the room temperature is 994.5 + / -46.29 MPa, and compared with a component manufactured through pre-alloying powder, the average UTS is improved by 14%, grains are refined by 14%, the geometric dislocation density is increased by 60.8%, and the cost is reduced by about 80%.

Description

Technical Field

[0001] The present invention relates to the field of Ti2AlNb alloy manufacturing, and particularly to a method for in-situ manufacturing of high-performance Ti2AlNb alloy by laser directed energy deposition with particle size control. Background Art

[0002] Ti2AlNb alloy is regarded as a potential alternative material to nickel-based superalloys (8 - 8.5 g / cm 3 ) in the aerospace field due to its excellent high-temperature performance and moderate density (5 - 5.7 g / cm 3 ). It can effectively reduce the structural weight and work stably in a high-temperature environment of 600 - 700 °C, with broad application prospects. The manufacturing of Ti2AlNb alloy mainly relies on casting, thermomechanical processing (such as forging and rolling), and powder metallurgy technology. However, these traditional methods have obvious limitations: due to the density and melting point differences of Ti, Al, and Nb elements in the casting process, it is easy to cause coarse grains, Al burning loss, and Nb segregation, thus affecting the mechanical properties; the difficult deformation characteristics and high-temperature brittleness of the alloy limit its thermomechanical processing applications. Although the development of powder metallurgy technology has brought new opportunities for the manufacturing of Ti2AlNb-based alloys, the expensive molds and high porosity hinder its applications.

[0003] Additive manufacturing has advantages such as customized forming, low cost, high material utilization rate, and short production cycle, and has broad application scenarios in the production of Ti2AlNb alloys. Due to the low plasticity of Ti2AlNb alloys at room temperature, it is difficult to make wires. Therefore, scholars mostly use the multi-wire feeding process for arc additive manufacturing of Ti2AlNb alloys. Li et al. used TiNb-Al dual-wire arc additive manufacturing to produce Ti2AlNb-based alloys. Due to composition and microstructure segregation, the UTS of the prepared alloy was only 504 MPa, 0.41. Yu et al. improved the uniformity of arc additive manufacturing of Ti2AlNb alloys by optimizing the droplet transfer mode, and the highest UTS was 891 MPa, but coarse columnar grains were prone to appear in the alloy. Polozov et al. used Ti, Al, and Nb elemental powders as raw materials and successfully prepared a high-density Ti2AlNb alloy through selective laser melting (SLM) combined with hot isostatic pressing (HIP) and subsequent heat treatment, but there were still un-melted Nb particles in the alloy. In another study, Polozov et al. performed high-temperature annealing on the SLM-in-situ-prepared Ti2AlNb alloy, which improved the problem of un-melted niobium particles, but the UTS of the alloy was only 286 MPa. Zhou et al. used pre-alloyed powders as raw materials and prepared high-performance Ti2AlNb alloys by SLM technology, with the maximum UTS of 1090 MPa. Polozov et al. used alloyed powders for SLM manufacturing of Ti2AlNb alloys, and the prepared alloys had cracks, with a UTS of 693 MPa. Zhou et al. prepared Ti2AlNb alloys by the SLM process, and the UTS after solution + aging treatment was 973 MPa. Li et al. used TiNb and Al wires as raw materials and innovatively adopted the dual-wire parallel feeding technology to successfully prepare Ti2AlNb alloys by electron beam freeform fabrication, with the UTS approaching 879.5 MPa. Although using alloyed welding wires or powders can improve the alloy properties, it will significantly increase the cost. Using elemental welding wires or powders for additive manufacturing of Ti2AlNb alloys is difficult to obtain uniform element distribution and good mechanical properties. Therefore, this study proposes a powder particle size regulation model to fabricate Ti2AlNb alloys with excellent properties. Summary of the Invention

[0004] Aiming at the deficiencies of the existing technology, the purpose of the present invention is to provide a method for in-situ manufacturing of high-performance Ti2AlNb alloys by laser directed energy deposition through particle size regulation.

[0005] To achieve the above purpose, the present invention provides the following technical solutions:

[0006] A method for in-situ manufacturing of high-performance Ti2AlNb alloys by laser directed energy deposition through particle size regulation, comprising the following steps:

[0007] S1: Screen Ti, Al, and Nb powders, and ensure that the purity of the Ti powder is 99.8%, the purity of the Al powder is 99.9%, and the purity of the Nb powder is 99.9%;

[0008] S2: After adjusting the powder particle sizes of the screened Ti, Al, and Nb powders, pre-mix them in a planetary ball mill, and after mixing evenly, place them in a vacuum drying oven at 120 °C and dry for 120 minutes;

[0009] S3: Use sandpaper to polish the surface of the substrate to remove the oxide layer, and then clean and dry it with alcohol and acetone to remove residual impurities;

[0010] S4: Place the dried substrate under the LDED in-situ manufacturing device, introduce the Ti, Al, and Nb mixed powder into the LDED in-situ manufacturing device through H2, deposit layer by layer with a laser power of 1500 W and a scanning speed of 5 mm / s, and apply equivalent volumetric heat flux and cooling time layer by layer to each deposited layer through multi-scale simulation to manufacture a high-performance Ti2AlNb alloy.

[0011] Preferably, in step S1, the screened Ti, Al, and Nb powders are spherical elemental powders.

[0012] Preferably, in step S2, the method for adjusting the powder particle sizes is as follows: taking the Nb powder as the standard, screen the Ti powder with a diameter of 25 - 57 μm and an average diameter of 41 μm, and the Al powder with a diameter of 18 - 42 μm and an average diameter of 30 μm through a molecular sieve.

[0013] Preferably, in step S2, the ratio of Ti, Al, and Nb powders entering the planetary ball mill is 46:11:43; the rotation speed of the planetary ball mill is 100 r / min, and the mixing time is 180 minutes.

[0014] Preferably, in step S3, the substrate is a TA15 plate with dimensions of 60 mm × 50 mm × 5 mm.

[0015] Preferably, in step S4, in the LDED in-situ manufacturing, the laser source uses a wavelength of 1064 nm, a spot diameter of 3.0 mm, and the laser system is IPGYLS - 10000.

[0016] Preferably, in step S4, the powder feeding rate is 5.6 g / min.

[0017] The multi-scale simulation method in step S4 includes a micro-scale laser scanning model and a macro-scale laser scanning model, and specifically includes the following steps:

[0018] S41: Establish a micro-scale laser scanning model composed of the substrate TA15 and the first layer of powder;

[0019] S42: In the micro-scale laser scanning model, adopt graded mesh division, use fine meshes near the deposition layer and coarse meshes in other areas to obtain the molten pool size and temperature history of the single-pass scanning trajectory points;

[0020] S43: Select the middle position of the first layer of powder to obtain the molten pool size and temperature history, and read the time for the powder to rise from the initial temperature to the highest temperature from the temperature history;

[0021] S44: Simplify the laser heat source into an equivalent volumetric heat source with a length equal to the spot diameter, a width equal to the melt channel width, and a height equal to the molten pool depth, establish a macro-scale laser scanning model for a single layer of powder, calculate the equivalent volumetric heat source based on the powder heating time and molten pool size obtained from the micro-scale laser scanning model, and calculate the single-layer cooling time through the single-layer actual deposition time, so as to control the thermal cycle of the entire component.

[0022] Preferably, the time for the powder to rise from the initial temperature to the highest temperature obtained in step 43 is 0.1 s, and the molten pool size is a molten pool length of 0.6 mm, a melt channel width of 2.8 mm, and a molten pool depth of 0.7 mm.

[0023] Preferably, the calculation formula for the equivalent volumetric heat source is:

[0024]

[0025] where A is the laser absorption rate, P is the laser power, ds is the laser spot diameter, dm is the melt channel width, and h is the molten pool depth;

[0026] The single-layer cooling time is the single-layer actual deposition time minus the time for the powder to rise from the initial temperature to the highest temperature. The single-layer actual deposition time is 8 s, so the single-layer cooling time is 7.9 s.

[0027] Compared with the prior art, the beneficial effects of the present invention are:

[0028] In the present invention, through the powder regulation model, the Ti, Al, and Nb powders are fully melted and mixed, reducing the phenomena of element enrichment and depletion; controlling the thermal cycle for each deposition layer, the phase compositions in the top and middle regions of the component are B2 and B2+α2+O phases respectively. The uneven heat conduction accelerates the cooling rate, resulting in an increase in the dislocation density; the Nb particles with a relatively high melting point serve as heterogeneous nucleation substrates, and the heat released by the negative mixing enthalpy changes the local heat flow direction in front of the solid / liquid interface, thereby changing the original grain growth direction and finally refining the grains.

[0029] The average ultimate tensile strength (UTS) of this in-situ manufactured Ti2AlNb alloy by LDED at room temperature is 994.5 ± 46.29 MPa. Compared with the components manufactured from pre-alloyed powder, the average UTS is increased by 14%, the grain size is refined by 14%, and the geometric dislocation density is increased by 60.8%. Its performance is comparable to that of traditional processes such as SPS and EBAM, and the cost is reduced by about 80%. Description of the Drawings

[0030] Figure 1 It is a schematic structural diagram of the (a) LDED device in this embodiment; (b) schematic diagram of sampling for longitudinal tensile test; (c) schematic diagram of hardness test positions; (d) schematic diagram of sampling for transverse tensile test; (e) schematic diagram of the positions for microstructure characterization.

[0031] Figure 2 It is the front view of the thermophysical properties of Ti2AlNb and TA15 in this embodiment.

[0032] Figure 3 It is the flow chart for powder particle size regulation in this embodiment.

[0033] Figure 4 It is the morphology and particle size diagram of Ti2AlNb pre-alloyed powder and Nb, Ti, Al elemental powders in this embodiment.

[0034] Figure 5 It is the schematic diagram of the multi-scale model in this embodiment: (a) schematic diagram of the micro-scale laser scanning model; (b) curve of temperature history; (c) schematic diagram of the melt pool size; (d) schematic diagram of equivalent heat source application; (e) schematic diagram of the thermal cycle simulation of the whole component.

[0035] Figure 6 It is the schematic diagram of the microstructure and XRD analysis of the samples prepared from two kinds of powders in this embodiment.

[0036] Figure 7 It is the schematic diagram of the TEM analysis of the samples prepared from two kinds of powders in this embodiment.

[0037] Figure 8 It is the schematic diagram of the SEM-EDS elemental surface scanning analysis of the samples prepared from two kinds of powders in this embodiment.

[0038] Figure 9 It is the schematic diagram of the EBSD analysis of the samples prepared from two kinds of powders in this embodiment.

[0039] Figure 10 It is the schematic diagram of the EBSD analysis of the c and g regions of the samples prepared from two kinds of powders in this embodiment.

[0040] Figure 11Schematic diagram for the mechanical property analysis of samples prepared from two kinds of powders in this embodiment.

[0041] Figure 12 Schematic diagram for the thermal cycle XRD analysis of Ti2AlNb alloy samples in this embodiment.

[0042] Figure 13 Schematic diagram for the verification analysis of the finite element model in this embodiment.

[0043] Figure 14 Schematic diagram for the preparation of Ti2AlNb alloy by LDED in this embodiment.

[0044] Figure 15 Schematic diagram for the phase evolution mechanism in this embodiment.

[0045] Figure 16 Test result graph for the cooling rate in this embodiment.

[0046] Figure 17 (a) Schematic diagram of the grain structure of Ti2AlNb alloy manufactured by LDED using pre-alloyed powder; (b) Schematic diagram of grain refinement induced by Nb particles; (c) Schematic diagram of the influence of mixing enthalpy on grain growth.

[0047] Figure 18 Schematic diagrams of the tensile fracture morphology and overall fracture of samples manufactured from two kinds of powders in this embodiment.

[0048] In the accompanying drawings: 1 - planetary ball mill, 2 - sieve, 3 - device body, 31 - coaxial nozzle, 4 - ABB robotic arm, 5 - laser device, 51 - laser beam, 6 - industrial computer, 7 - powder feeder, 8 - powder distributor, 9 - working platform, 10 - manufacturing platform, 11 - TA15 plate, 12 - argon chamber, 13 - deposition layer. Specific embodiments

[0049] In order to make the objectives, technical solutions and advantages of the present invention clearer and more understandable, the present 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 only used to explain the present invention and are not used to limit the present invention. In addition, the technical features involved in the various embodiments of the present invention described below can be combined with each other as long as they do not conflict with each other.

[0050] Please refer to Figure 1-15 , a method for in-situ manufacturing of high-performance Ti2AlNb alloy by laser directed energy deposition through particle size regulation provided by the present invention, includes the following steps:

[0051] S1: Screen Ti, Al, and Nb powders. The Ti, Al, and Nb powders are spherical elemental powders, and ensure that the purity of the Ti powder is 99.8%, the purity of the Al powder is 99.9%, and the purity of the Nb powder is 99.9%.

[0052] S2: After adjusting the powder particle sizes of the screened Ti, Al, and Nb powders, perform pre-mixing using a planetary ball mill 1. After mixing evenly, place it in a vacuum drying oven at 120 °C and dry for 120 minutes.

[0053] Previous studies usually simply mixed powders of different particle sizes and used SLM to in-situ fabricate Ti2AlNb alloys. However, it was difficult to obtain a Ti2AlNb alloy with uniform composition, and its mechanical properties were poor. This powder particle size adjustment method can ensure that the mass ratio of single Ti, Al, and Nb powders is close to the theoretical mass ratio of each element in the Ti2AlNb alloy (Ti:Al:Nb = 46:11:43). The powder particle size adjustment process is as Figure 3 shown. After determining the average particle sizes of the Ti, Al, and Nb powder particles using the powder particle size adjustment model ( Figure 3 (a)), select the corresponding specification sieve 2 ( Figure 3 (b)). Taking the Nb powder as the standard, screen the Ti powder with a diameter of 25 - 57 μm and an average diameter of 41 μm, and the Al powder with a diameter of 18 - 42 μm and an average diameter of 30 μm through a molecular sieve; finally, put them into the planetary ball mill 1 for mixing. The rotation speed of the planetary ball mill 1 is 100 r / min, and the mixing time is 180 minutes ( Figure 3 (c)).

[0054] Specifically, the mass of a single spherical powder is:

[0055]

[0056] In the formula, d is the average particle size of a single spherical powder. The average particle sizes of the three powders satisfy the following relationship:

[0057]

[0058] ...... In the formula, m TTi , m ... TAl , m ... TNb are the theoretical masses of Ti, Al, and Nb in the Ti2AlNb alloy, ρ Ti , ρ Al , ρ Nb are the densities of Ti, Al, and Nb, and d Ti , d Al , d Nb are the average particle sizes of the Ti, Al, and Nb elemental powders.

[0059] Since Nb powder is relatively expensive, taking the particle size of the purchased Nb powder (average value is 31.99um) as the standard, the average particle sizes of Ti and Al are calculated to be 40.54um and 29.84um through formula (2). According to the average particle sizes of Ti and Al powders, the aperture ranges of sieve 2 are selected, which are 25 - 57um (average aperture is 41um) and 18 - 42um (average aperture is 30um) respectively. The finally sieved powder particle sizes are 41.52um and 30.57um respectively, which are close to the calculated values. The morphologies and particle size distributions of Ti2AlNb pre-alloyed powder and Nb, Ti, Al elemental powders are as Figure 4 (a)-(h) shown, where Figure 4 (a, b, c, d) are the SEM morphologies of powders Ti2AlNb (a), Nb (b), Ti (c), Al (d); Figure 4 (e, f, g, h) are the particle size distributions of Ti2AlNb (e), Nb (f), Ti (g), Al (h); Figure 4 (i, j, k, l) are the EDS results of the mixed powder particles: mixed image (i), Ti (j), Al (k) and Nb (l). The results show that the three powder particles of Ti, Al, and Nb are evenly distributed after mixing, and no obvious segregation or local absence phenomenon is observed.

[0060] S3: Use sandpaper to polish the surface of the substrate to remove the oxide layer, and then clean and dry it with alcohol and acetone to remove residual impurities. The substrate is a TA15 plate with dimensions of 60mm × 50mm × 5mm;

[0061] S4: Place the dried substrate under the LDED in-situ manufacturing device. The laser source has a wavelength of 1064nm, a spot diameter of 3.0mm, and the laser system is IPGYLS - 10000. Bring the Ti, Al, Nb mixed powder into the LDED in-situ manufacturing device through H2. The powder feeding rate is 5.6g / min. Deposit layer by layer at a laser power of 1500W and a scanning rate of 5mm / s. Apply equivalent volumetric heat flux and cooling time to each deposited layer to manufacture high-performance Ti2AlNb alloy. Finally, two 70-layer Ti2AlNb alloy thin-walled walls are additively manufactured, and the surface is observed to have a metallic luster and no cracks are generated.

[0062] As Figure 1As shown in (a), the LDED in-situ manufacturing device is located inside the argon chamber 12, and specifically includes a device body 3, an ABB robotic arm 4, a laser device 5, an industrial computer 6, a powder feeder 7, and a powder distributor 8. A coaxial nozzle 31 is provided on the device body 3. The ABB robotic arm 4 controls the device body 3 to be above the working platform 9. A manufacturing platform 10 is placed on the working platform 9. A TA15 plate 11 is sequentially placed on the manufacturing platform 10. The industrial computer 6 controls the laser device 5 to emit a laser beam 51 to the device body 3. At the same time, through Ar, the Ti, Al, and Nb mixed powder in the powder feeder 7 is transported to the powder distributor 8. After being distributed, it is input into the coaxial nozzle 31 of the device body 3 and ejected, forming a deposition layer 13 of the Ti2AlNb alloy sample on the TA15 plate 11.

[0063] The multi-scale simulation method in step S4 is used to efficiently calculate the LDED thermal cycle process of the Ti2AlNb alloy, including a micro-scale laser scanning model and a macro-scale laser scanning model. As Figure 5 shown, through the micro-scale laser scanning model, the molten pool size and temperature history of the points on the single-track scanning trajectory are obtained, so as to calculate the equivalent volumetric heat flux. In the macro-scale laser scanning model, the equivalent volumetric heat flux and cooling time are applied layer by layer, thus completing the thermal cycle simulation of the entire component.

[0064] Specifically, it includes the following steps:

[0065] S41: Establish a micro-scale laser scanning model composed of the substrate TA15 and the first layer of powder, as Figure 5 (a) shown;

[0066] S42: In the micro-scale laser scanning model, use a graded grid division, with a fine grid near the deposition layer and a coarse grid in other areas to obtain the molten pool size and temperature history of the points on the single-track scanning trajectory;

[0067] In the micro-scale laser scanning model, the DC3D8 linear heat transfer hexahedron element is used, and the heat source model is a double ellipsoid heat source, and its formula is:

[0068]

[0069] Among them, P is the laser power, a f 、a r respectively represent the semi-axes lengths of the ellipsoid in the scanning direction of the front and rear edges of the laser beam, a f = 3mm, a r = 2mm, b and c respectively represent the semi-axes lengths of the ellipsoid perpendicular to the scanning direction and the depth direction. It is proved that b = 2.5mm, c = 0.4mm, f1 and f2 represent the energy distribution coefficients before and after the heat source, f1 + f2 = 2, and f1 = 0.4, f2 = 1.6 are taken.

[0070] S43: Select the middle position of the first layer of powder to obtain the melt pool size and temperature history, and read the time for the powder to rise from the initial temperature to the highest temperature in the temperature history;

[0071] During the deposition process, the laser parameters are constant and the melt pool size fluctuates little. Therefore, select the middle position of the first layer of powder to obtain the melt pool size and temperature history. The results are as Figure 5 (b) and (c). Under the action of the laser, the powder experiences instantaneous heating and cooling. The time for the obtained powder to rise from the initial temperature to the highest temperature is 0.1 s, and the melt pool size is a melt pool length of 0.6 mm, a melt channel width of 2.8 mm, and a melt pool depth of 0.7 mm.

[0072] S44: Simplify the laser heat source into an equivalent volumetric heat source with a length of the spot diameter, a width of the melt channel width, and a height of the melt pool depth. Establish a macro-scale laser scanning model for a single layer of powder. Calculate the equivalent volumetric heat source based on the powder heating time and melt pool size obtained from the micro-scale laser scanning model, and calculate the single-layer cooling time through the single-layer actual deposition time, so as to control the thermal cycle of the entire component.

[0073] The grid processing method of the macro-scale laser scanning model is the same as that of the micro-scale laser scanning model. In the macro-scale laser scanning model, the nodes experience a heating and cooling process similar to that of the micro-scale laser scanning model. Therefore, it is reasonable to apply the temperature history of a single node in the micro-scale laser scanning model to the macro-scale laser scanning model. In the macro-scale laser scanning model, the heat source is simplified into an equivalent volumetric heat source with a length of the spot diameter, a width of the melt channel width, and a height of the melt pool depth. The calculation formula is as shown in Equation (4).

[0074] The calculation formula for the equivalent volumetric heat source is:

[0075]

[0076] where A is the laser absorption rate (A = 37.72%), P is the laser power, ds is the laser spot diameter, dm is the melt channel width, and h is the melt pool depth;

[0077] The single-layer cooling time is the single-layer actual deposition time minus the time for the powder to rise from the initial temperature to the highest temperature. Calculate the equivalent volumetric heat source based on the powder heating time (0.1 s) and melt pool size obtained from the micro-scale laser scanning model. This heat source is used to heat a single layer of powder, as Figure 5 (d) shows. Since the single-layer actual deposition time is 8 s, the single-layer cooling time is set to 7.9 s. Finally, apply the equivalent volumetric heat flux and cooling time layer by layer to each deposition layer to complete the thermal cycle control of the entire component. The results are as Figure 5 (e).

[0078] The performance of the in-situ fabricated Ti2AlNb alloy by LDED is verified through experimental analysis as follows.

[0079] I. Tensile test, hardness test and microstructure characterization of Ti2AlNb alloy samples

[0080] Two 70-layer Ti2AlNb alloy thin-walled walls were cut into three samples for testing the tensile properties in the longitudinal direction ( Figure 1 (b)) and transverse direction ( Figure 1 (d)). The tensile experiment was completed on a KPL-Focal U100 (China) device. The experimental method referred to the Chinese standard GB / T 228.1-2010 "Tensile test at room temperature for metallic materials". The strain rate was 1×10 -3 s -1 , the tensile rate was 0.5 mm / min, and a video extensometer was used to measure the strain. The final UTS was the average value of the three samples.

[0081] Five positions were taken at equal intervals (6 mm) from the bottom to the top of the cut samples, and 3 points were tested in the horizontal direction at each position, with a total of 15 points tested. The hardness test positions are as shown in Figure 1 (c).

[0082] Samples from the middle and top of the Ti2AlNb alloy thin-walled wall were taken for grinding and polishing, and then microstructure characterization was carried out, as shown in Figure 1 (e). After the samples were corroded with Kroll reagent (HF:HNO3:H2O = 1:6:10), the microstructure was observed using a scanning electron microscope (SEM, ZEISS GeminiSEM 300, Germany), and the elemental distribution of the samples was analyzed by EDS. The phase composition was determined using X-ray diffraction (XRD, Rigaku Ultima IV, Japan). The EBSD characterization of the samples prepared from the pre-mixed powder and pre-alloyed powder was carried out using an EDA-TSL (USA) device. During the test, the acceleration voltage was 15 kV, the beam current was 1 nA, and the working distance was 14 mm. Among them, the grain size and kernel average misorientation (KAM) analysis were scanned with a step size of 1 μm, and the O / α2 phase precipitation characteristics analysis was scanned with a fine step size of 0.1 μm. The EBSD data was analyzed using AZtecCrystal software, and the samples were prepared by vibratory polishing method. And the phase of the microstructure was analyzed using a transmission electron microscope (TEM, JEM-2100F, Japan).

[0083] II. Thermophysical properties of Ti2AlNb alloy samples

[0084] The thermophysical properties of Ti2AlNb and TA15 were calculated by JMatPro software, as shown in Figure 2 .

[0085] III. Thermal Cycle XRD Analysis of Ti2AlNb Alloy Samples

[0086] XRD analysis shows that region b of the deposited sample only contains the B2 phase, while region c contains the B2, O, and α2 phases. To explore the phase evolution mechanism of the deposited sample, it is necessary to further analyze the thermal cycle during deposition. Points A (located in region b) and B (located in region c) in the finite element model are selected for analysis, as shown in Figure 12 (a). Combining with the pseudo-binary phase diagram of the Ti-22Al-xNb alloy (see Figure 12 (b)), when the temperature is higher than 1060 °C, the alloy phase composition is a single B2 phase; when the temperature is lower than 1060 °C, the α2 and O phases start to precipitate. Therefore, taking 1060 °C as the boundary, the thermal cycle with the lowest temperature greater than 1060 °C is defined as the high-temperature thermal cycle (HTTC); the thermal cycle with the highest temperature less than 1060 °C is defined as the low-temperature thermal cycle (LTTC); if 1060 °C is between the lowest and highest temperatures, it is defined as the medium-temperature thermal cycle (MTTC). It can be seen that point A has experienced 10 HTTCs, point B has experienced 6 MTTCs, and 29 LTTCs. As the deposition progresses, the temperature cumulative effect causes the peak temperature at point B (2477 °C) to point A (2801 °C) to gradually increase, as shown in Figure 12 (c).

[0087] An infrared camera (FLIRA615) is used for temperature field measurement. Before the experiment, the infrared data is calibrated using a K-type thermocouple (KPS-IN600-K, 10Hz). The thermocouple is placed on the left side of the Ta15 substrate near the additive manufacturing starting position( Figure 13 (a)). The calibration results show that the trend of the thermocouple data is in good agreement with the infrared temperature measurement data( Figure 13 (b)), verifying the reliability of the infrared temperature measurement method. To verify the accuracy of the model, the surface temperature of the deposited sample is measured using an infrared camera. The positions of the selected points are shown in Figure 13 (c), corresponding to point A in the finite element model. The measured temperature is compared with the simulation results, and the results are shown in Figure 13 (d). Point A has experienced HTTC. The error between the simulated peak temperature (2658 °C) and the measured temperature (2768 °C) is 3.9%. Although there is a certain error between the simulated data and the measured data, the temperature change trends of the simulation and the actual measurement are basically the same.

[0088] IV. Microstructure and Element Distribution of Deposited Samples of Premixed Powder and Prealloyed Powder

[0089] (1) Microstructure and XRD Analysis of Premixed Powder and Prealloyed Powder

[0090] As shown in Figure 6As shown in Figs. (a) and 6(d), representative regions were selected at the top (regions b, f), middle (regions c, g), and bottom (regions d, h) of the samples prepared from the pre-mixed powder and pre-alloyed powder, respectively, for microstructure and XRD analysis.

[0091] The phase composition of the Ti2AlNb alloy varies under different processing techniques, mainly including three basic phases: the B2 / β phase (body-centered cubic structure, Pm3m symmetry), the O phase based on Ti2AlNb (ordered orthorhombic structure, Cmcm symmetry), and the α2 phase based on Ti3Al (ordered close-packed hexagonal structure, P63 / mmc symmetry). In addition, under SEM observation, the colors of the B2 / β, O, and α2 phases are light, gray, and dark, respectively. In the Ti2AlNb alloy, the B2 phase is the plastic phase, while the O and α2 phases exist as strengthening phases.

[0092] From this microstructure and XRD analysis, as Figure 6 Figs. (b)-(d) and 6(f)-(h) show the microstructures of the samples prepared from the pre-mixed powder and pre-alloyed powder, respectively. Combining with the XRD Figure 6 (i)-(k)) analysis reveals that: the top regions (b, f) are composed only of the B2 phase, and the grain boundaries are clearly visible; while the middle (c, g) and bottom (d, h) regions are composed of the B2 phase, O phase, and α2 phase. The microstructures are similar, with the O and α2 phases precipitating in the B2 matrix, showing typical microstructural features. The samples obtained by the two preparation methods show the same tissue evolution law along the deposition height direction: the tissue from the bottom to the top gradually transitions from the B2+O+α2 three-phase to the single B2 phase.

[0093] (2) TEM analysis of the pre-mixed powder and pre-alloyed powder

[0094] To more clearly distinguish the O phase from the α2 phase, TEM analysis was performed on the needle-like precipitates in region c.

[0095] Figure 6 Fig. (a) shows the bright-field image of this region. Three characteristic regions (A, B, C) were selected for selected area electron diffraction (SAED), and the results are as Figure 6 Figs. (b)-(c), corresponding to the B2 matrix phase, O phase, and α2 phase, respectively. The O and α2 phases are both needle-shaped, but the O phase has a relatively smaller size, and the phase composition is consistent with the XRD test results. In addition, by performing inverse Fourier transform (IFFT) on the high-resolution image Figure 7 (e) near the grain boundary of the α2 phase (region D) (see Figure 7(f)), it is found that there are a large number of dislocations and lattice distortions near the grain boundaries. The interaction and interlacing of dislocations in different directions make it difficult for them to move, achieving the effect of dislocation pinning, and ultimately effectively improving the tensile strength of the Ti2AlNb alloy.

[0096] At present, there are very few studies on the in-situ additive manufacturing of Ti2AlNb alloy using pre-mixed powders of Ti, Al, and Nb as raw materials. Only the Polozov team has made some attempts. Although they successfully fabricated Ti2AlNb alloy by in-situ SLM, there are problems of unmelted Nb particles and uneven composition distribution. Although the uneven composition distribution can be improved by subsequent heat treatment, the mechanical properties of the alloy are poor.

[0097] (3) SEM-EDS elemental surface scanning analysis of pre-mixed powders and pre-alloyed powders

[0098] Through SEM-EDS elemental surface scanning analysis of the elemental distribution around the molten pool of the pre-mixed powder sample ( Figure 8 (a)-(d)), the results show that the Nb element has been completely melted inside the molten pool, and no residual unmelted Nb particles are observed. Further, high-magnification SEM-EDS analysis is carried out on the pre-mixed powder region c ( Figure 8 (e)-(h)) and the elemental distribution in the middle part of the sample prepared in the pre-alloyed powder region g ( Figure 8 (i)-(l)). The results show that only a small amount of segregation occurs at the grain boundary positions in the two specimens, and there is no macroscopic segregation in general. The Ti, Al, and Nb elements in the sample prepared from the pre-mixed powder are evenly distributed. The measured mass ratio (47.4:10.3:42.3) has a small deviation from the theoretical mass ratio (46:11:43), and no unmelted Nb particles are detected at the molten pool boundary, indicating that this method effectively solves the problem of uneven composition distribution in the in-situ additive manufacturing of Ti2AlNb alloy with pre-mixed powders.

[0099] (4) EBSD analysis of pre-mixed powders and pre-alloyed powders

[0100] The EBSD analysis results of the middle regions ( Figure 9 regions c and g) of the Ti2AlNb alloy samples fabricated from pre-mixed powders and pre-alloyed powders. The average grain sizes are 211.3 μm and 245.7 μm respectively ( Figure 9 (b) and (f)). The grains of the sample fabricated from the pre-mixed powder are refined by 14%, and the grain sizes are all less than 400 μm, while some grain sizes in the sample fabricated from the pre-alloyed powder are greater than 400 μm ( Figure 9 (a)). At the same time, the average misorientation (Kernel Average Misorientation, KAM) of both is analyzed, and the results are as Figure 9(c) and (g) show that the KAM (1.39°) of the samples made of pre-mixed powder is significantly higher than that of the pre-alloyed powder samples (0.54°), and the average misorientation increases by 61.2% ( Figure 9 (e)). The higher KAM value indicates a higher dislocation density (GND) in the samples. Figure 9 (i) and (k) are the pole figures of the two. The texture of the samples made of pre-mixed powder is relatively weakened compared with that of the samples made of pre-alloyed powder, and the preferred grain orientation is weakened. The B2 phase is the main plastic phase in the Ti2AlNb alloy and has more slip systems. Among them, {110}<111> and {112}<111> are the main slip systems. The Schmid factors of the pre-mixed powder and pre-alloyed powder samples on these two slip systems are calculated: The results are as Figure 9 (j) and (l). Whether on the {110}<111> or {112}<111> slip system, the Schmid factor (0.45) of the pre-mixed powder sample is less than that of the pre-alloyed powder sample (0.47), indicating that it is more difficult to activate the slip system of the pre-mixed powder sample, which may be the reason for the increase in tensile strength.

[0101] In addition, high-resolution EBSD (step size 0.1 μm) was used to comparatively analyze the microstructure, precipitate orientation and its volume fraction in the pre-mixed powder (region c) and pre-alloyed powder (region g) samples. The results show that the microstructure of region c is thinner than that of region g ( Figure 10 (a), (d)). The precipitates in both samples show significant intragranular orientation characteristics ( Figure 10 (b), (e)). Pole figure analysis ( Figure 10 (i)) found that the B2, O and α2 phases of the pre-mixed powder sample (region c) strictly follow the orientation relationships of (001)O / / {0001}α2 and (001)O / / {110}B2, which conform to the classical Burgers orientation relationship. Figure 10 (c) and (f) are the phase composition diagrams of region c and region g, where the B2 phase is green, the O phase is purple, and the α2 phase is yellow. Through statistics by AZtecCrystal software, the contents of the B2 phase, O phase and α2 phase in the pre-mixed (region c) powder sample are 72.7%, 9.1% and 18.2% respectively; while the contents of the B2 phase, O phase and α2 phase in the pre-alloyed (region f) powder sample are 71.6%, 25.2% and 3.1% respectively, as Figure 10(h). It is worth noting that although the α2 phase in both the pre-mixed (region c) and pre-alloyed (region g) samples exhibits a nearly needle-like morphology, there are obvious differences in the morphology of the O phase between the two: the O phase in the pre-mixed sample is nearly needle-like, while the O phase in the pre-alloyed sample is a lamellar structure. A faster cooling rate will cause the lamellar structure of the O phase to become thinner and the volume fraction to decrease, and induce the precipitation of needle-like O phase in the matrix. Since the cooling rate of the pre-mixed powder sample is higher than that of the pre-alloyed powder, its O phase volume fraction is lower, the microstructure is thinner, and needle-like precipitates are formed in the B2 matrix. This observation is completely consistent with the EBSD analysis.

[0102] (5) Mechanical property analysis of pre-mixed powder and pre-alloyed powder

[0103] The tensile test results show that the longitudinal and transverse average tensile strengths (UTS) and average elongation rates (EL) of the in-situ fabricated samples of pre-mixed powder are 973.75 ± 36.85 MPa, 1.07 ± 0.1% and 1016 ± 55.72 MPa, 1.11% ± 0.2% ( Figure 11 (a) and (c)); while the average UTS and average EL of the samples fabricated from pre-alloyed powder are 851.24 ± 20.75 MPa, 0.85 ± 0.03% and 849 ± 43.2 MPa, 0.89% ± 0.1% ( Figure 11 (b) and (d)). The results show that the average UTS and EL of the pre-mixed powder samples are both better than those of the pre-alloyed powder samples.

[0104] In addition, the hardness test results along the deposition direction (5 positions at equal intervals of 6 mm, 3 test points in the horizontal direction at each position) ( Figure 11 (e)) show that the average hardness of the pre-mixed powder and pre-alloyed powder samples in the B2 + O + α2 phase region are 436 HV and 424 HV respectively, and the average hardness in the B2 phase region is 363 HV and 361 HV respectively. The hardness distribution trends of the samples prepared from the two powders are basically the same. The hardness in the B2 + O + α2 region is higher than that in the B2 phase region, which is attributed to the existence of the α2 / O phase as a strengthening phase and the B2 phase as a plastic phase in the Ti2AlNb alloy. However, the hardness values in the horizontal and vertical directions in the same phase composition region have little difference, confirming that the overall performance of the pre-mixed powder sample is uniform.

[0105] To evaluate the position of the LDED in-situ manufacturing process among common processes, we compared its mechanical properties with those of components manufactured by common processes, such as Figure 11As shown in (f). The results show that the UTS of the components in-situ fabricated by LDED is close to that of SLM and SPS (Spark Plasma Sintering), and is better than that of PM (Powder Metallurgy) and WAAM (Wire Arc Additive Manufacturing). Currently, the process with better performance for in-situ fabricating Ti2AlNb alloy is EBAM (Electron Beam Additive Manufacturing). Although Polozov et al. also fabricated Ti2AlNb alloy in-situ using SLM, the mechanical properties are poor. The mechanical properties of the components fabricated by the LDED in-situ manufacturing process proposed in the present invention are close to those of EBAM, and the cost is lower. However, compared with the components fabricated by forging and some SLM processes, there is still a gap in mechanical properties, especially in terms of plasticity.

[0106] V. Analysis of Alloy Element Homogenization and Phase Transformation Mechanism

[0107] (1) Homogenization Analysis of Three Kinds of Powders

[0108] After mixing the three kinds of powders evenly and feeding them into the molten pool, the powders are randomly distributed. If the powder particle sizes are not optimized, it is easy to produce uneven element distribution. Taking the Nb element as an example for analysis, as Figure 14 shown in (a), the particle size of the Nb powder in Region 1 is small (with a small mass), and it can be fully melted, but the Nb content is lower than the alloy composition. Due to the fast cooling rate of the molten pool, it is difficult for the Nb elements in other regions to flow to Region 1, resulting in the generation of Nb-depleted regions. The particle size of the Nb powder in Region 2 is large (with a large mass), and un-melted Nb powder is likely to be produced. Even if it is fully melted, the remaining Nb elements do not have time to flow to other regions, and finally a Nb-rich region is formed. This phenomenon is consistent with what was observed by Polozov, Grigoriev, and Wang et al.

[0109] In the present invention, the particle sizes of Ti, Al, and Nb powder particles are optimized, so that the mass ratio of single Ti, Al, and Nb powders is close to the theoretical mass ratio of each element in the Ti2AlNb alloy (Ti:Al:Nb = 46:11:43). The average particle size values of the three kinds of powders are 41.52um, 30.57um, and 31.99um respectively. The particle size of the Nb powder is moderate and it is relatively easy to melt. As Figure 14 shown in (b), when three adjacent Ti, Al, and Nb powders that meet the particle size ratio are melted together, the in-situ generated Ti2AlNb alloy meets the theoretical mass ratio, and there will be no lack or surplus of Nb elements, thus reducing the occurrence probability of element deficiency or enrichment phenomena.

[0110] (2) Phase Transformation Mechanism Analysis of Three Kinds of Powders

[0111] As Figure 15 shown, when the laser acts on the pre-mixed powders, the powders are quickly melted, the temperature quickly rises to the peak, and then at a rate exceeding 103 The cooling rate in K / s drops sharply, and the α2 and O phases cannot precipitate in time before the B2 phase forms, resulting in a B2-phase composition for the Ti2AlNb alloy that has not undergone thermal cycling, as shown in Figure 15 (a). As shown in Figure 12 (c), point A is completely under HTTC conditions, and the O and α2 phases do not precipitate, so the phase composition is a single B2 phase. Point B has undergone 5 MTTCs and 35 LTTCS. During MTTC, the precipitation and dissolution of the α2 phase alternate, and the dissolution temperature is higher than the precipitation temperature, resulting in insignificant precipitation of the α2 phase. After entering the LTTC stage, when the thermal cycling temperature range is 1010 °C - 1060 °C, the α2 phase begins to precipitate along the grain boundaries ( Figure 15 (b)), and when the temperature is lower than 1010 °C, the plate-shaped O phase begins to precipitate ( Figure 15 (c)), as shown in Figure 13 (c), ultimately resulting in a phase composition of B2, O, and α2 phases at point B. In addition, the formation mechanism of the O phase in Ti2AlNb alloy components fabricated from pre-mixed powders is described. In the present invention, the precipitation sequence of the O phase is as follows: First, the pre-mixed Ti, Al, and Nb particles melt under the action of a laser to form a liquid phase. Then, during solidification, the B2 phase nucleates from the liquid phase. After LTTC, the O phase begins to precipitate, and an intermediate transition phase B19 appears during the precipitation process, i.e., the B2→B19→O phase transformation occurs.

[0112] An infrared camera was used to obtain the temperature curves at the midpoints (points A1, A2) of the first layer and the midpoints (points B1, B2) of the 35th layer of samples fabricated from pre-mixed powders ( Figure 16 (a)) and pre-alloyed powders ( Figure 16 (b)), and the corresponding cooling rates were calculated. The results are shown in Figure 16 (c) and (d): The cooling rates of the first layer of samples fabricated from pre-mixed powders and pre-alloyed powders are 1563 °C / s and 898 °C / s, respectively; the cooling rates of the 35th layer are 934 °C / s and 795 °C / s, respectively. It can be found that whether at the midpoint of the first layer or the 35th layer, the cooling rate of the pre-mixed powder is always higher than that of the pre-alloyed powder. A faster cooling rate is beneficial for forming a fine grain structure, which not only increases the number of grain boundaries but also causes more significant lattice distortion. Since the grain orientations on both sides of the grain boundary are different, dislocations generated by slip in one grain cannot be directly transmitted to the adjacent grain [3], resulting in dislocation pile-up and a significant increase in dislocation density, thereby leading to a higher KAM value.

[0113] The grain morphology of components fabricated from pre-alloyed powders is mainly columnar and equiaxed grains, as shown in Figure 17 (a); Figure 17(b) is a schematic diagram of heterogeneous nucleation. Under the action of laser, the melting time of Nb particles is later than that of Ti and Al particles. Nb can serve as an effective nucleation substrate for the B2 phase, promoting nucleation and refining the microstructure. The liquid mixing enthalpy of the pre-mixed powder of Ti, Al, and Nb is less than zero, indicating that the alloying process is an exothermic reaction. The released heat will change the direction of the heat flow in the local area in front of the solid / liquid interface. When depositing the next layer, the epitaxial growth of the original grains along the deposition direction is interrupted, thereby promoting the nucleation and growth of new grains in front of the solid / liquid interface, resulting in microstructure refinement, as Figure 17 shown in (c).

[0114] The pre-mixed powder sample has a higher KAM value than the pre-alloyed powder sample, which indicates that there is a greater dislocation density and more significant dislocation accumulation in the pre-mixed powder sample. This dislocation accumulation promotes the rearrangement process of dislocations, generating more dislocation boundaries through dislocation splitting or the formation of new dislocation walls, which are then transformed into low-angle grain boundaries (LAGB). Figure 9 (d) and (h) clearly show that the density of low-angle grain boundaries (2 - 15°) inside the grains of the pre-mixed powder sample is significantly higher than that of the pre-alloyed powder sample. During the subsequent dynamic recrystallization process, these low-angle grain boundaries will gradually evolve into high-angle grain boundaries (HAGB > 15°), ultimately promoting grain refinement.

[0115] The tensile fracture surfaces of the pre-mixed powder and pre-alloyed powder samples are as Figure 18 (a) and (b). The fracture of the pre-mixed powder specimen shows a mixed-mode characteristic: there are simultaneously intergranular fracture, transgranular fracture, and a small amount (about 7%) of quasi-cleavage fracture. It should be noted that the toughening effect of the small amount of cleavage fracture is not sufficient to offset the brittle effect brought by the coarse grains, so it still shows brittle fracture as a whole. The pre-alloyed powder sample shows typical brittle fracture characteristics, mainly manifested as intergranular fracture and transgranular fracture modes. In the intergranular fracture region ( Figure 18 (c) and (e)), the rock-like pattern structure formed by coarse equiaxed grains can be clearly observed, and obvious cracks exist between the grains. In the transgranular fracture region ( Figure 18 (a) and (f)), a large number of typical river patterns can be seen, and their directions are perpendicular to the grain boundaries. The quasi-cleavage fracture morphology observed in the tensile fracture surface of the pre-mixed powder (as Figure 18 shown in (d)) is consistent with these characteristics, confirming that its fracture surface shows quasi-cleavage characteristics rather than dimpled structure. This fracture surface morphology is highly consistent with the research results reported in the literature, all showing typical brittle fracture characteristics.

[0116] In summary, under the same conditions, compared with pre-alloyed powders, the advantages of using pre-mixed powders for in-situ manufacturing of Ti2AlNb alloys by LDED are mainly in terms of performance and cost: compared with samples manufactured with pre-alloyed powders, the average UTS increased by 12.5% and the average EL increased by 20.6%, and the performance is comparable to that of Ti2AlNb alloys manufactured by traditional processes. Statistical analysis shows that the average price for purchasing 1 kg of Ti2AlNb pre-alloyed powder is $778.18, while the average price for purchasing 1 kg of pre-mixed powder (Ti: 460 g, Al: 110 g, Nb: 430 g) is $155.64, resulting in an 80% cost reduction. It can be seen that using pre-mixed powders can manufacture high-performance Ti2AlNb alloy components at low cost, which is conducive to rapidly promoting the commercial application of Ti2AlNb alloys.

[0117] The present invention uses Ti, Al, and Nb elemental powders as raw materials to in-situ manufacture Ti2AlNb alloys with excellent performance by LDED:

[0118] (1) Through the powder regulation model, the mass ratio of single Ti, Al, and Nb powders is made close to the theoretical mass ratio of each element in the Ti2AlNb alloy (Ti:Al:Nb = 46:11:43). When three adjacent Ti, Al, and Nb powders that meet the particle size ratio are melted together, the in-situ generated Ti2AlNb alloy meets the theoretical mass ratio, without lacking or remaining elements, thereby reducing the probability of element deficiency or enrichment. The deviation between the mass of Ti, Al, and Nb in the deposited Ti2AlNb alloy component (47.4:10.3:42.3) and the theoretical mass is small, the element distribution is uniform, and no obvious macroscopic segregation phenomenon appears.

[0119] (2) Only the HTTC is experienced in the top region of the component, and the phase composition is a single B2 phase; the LTTC is dominant in the middle region, and the phase composition is B2 + α2 + O phase.

[0120] (3) Nb particles with a higher melting point serve as heterogeneous nucleation substrates, and the heat released by the negative mixing enthalpy changes the local heat flow direction in front of the solid / liquid interface, changing the growth direction of the original grains, thereby refining the grains.

[0121] (4) Due to grain refinement and dislocation strengthening, the average UTS and EL of the Ti2AlNb alloy manufactured by in-situ LDED using pre-mixed powders are 973.75 ± 36.85 MPa and 1.07 ± 0.1%, respectively, exceeding the components manufactured with pre-alloyed powders under the same process, and the performance is comparable to that of the SPS and EBAM processes.

[0122] The above are only specific embodiments 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 in-situ manufacturing of high-performance Ti2AlNb alloy by laser directed energy deposition with particle size control, characterized in that, It includes the following steps: S1: Screen Ti, Al, and Nb powders, and ensure that the purity of the Ti powder is 99.8%, the purity of the Al powder is 99.9%, and the purity of the Nb powder is 99.9%; S2: After adjusting the powder particle size of the screened Ti, Al, and Nb powders, perform pre-mixing in a planetary ball mill, and after mixing evenly, place them in a vacuum drying oven at 120 °C and dry for 120 minutes; S3: Use sandpaper to polish the surface of the substrate to remove the oxide layer, and then clean and dry it with alcohol and acetone to remove residual impurities; S4: Place the dried substrate under the LDED in-situ manufacturing device, introduce the Ti, Al, and Nb mixed powders into the LDED in-situ manufacturing device through H2, deposit layer by layer with a laser power of 1500 W and a scanning rate of 5 mm / s, and apply equivalent volumetric heat flux and cooling time layer by layer to each deposited layer through multi-scale simulation to manufacture high-performance Ti2AlNb alloy.

2. A method for in-situ fabricating a high-performance Ti2AlNb alloy by laser directed energy deposition through particle size control according to claim 1, characterized in that In step S1, the screened Ti, Al, and Nb powders are spherical elemental powders.

3. A method for in-situ manufacturing of high-performance Ti2AlNb alloy by laser directed energy deposition through particle size regulation according to claim 1, characterized in that, In step S2, the method for adjusting the powder particle size is: taking the Nb powder as the standard, screening the Ti powder with a molecular sieve to have a diameter of 25 - 57 μm and an average diameter of 41 μm, and the Al powder to have a diameter of 18 - 42 μm and an average diameter of 30 μm.

4. A method for in-situ fabricating a high-performance Ti2AlNb alloy by laser directed energy deposition through particle size control according to claim 3, characterized in that, In step S2, the ratio of Ti, Al, and Nb powders entering the planetary ball mill is 46:11:43; the rotation speed of the planetary ball mill is 100 r / min, and the mixing time is 180 minutes.

5. A method for in-situ manufacturing of high-performance Ti2AlNb alloy by laser directed energy deposition with particle size regulation according to claim 1, characterized in that, In step S3, the substrate is a TA15 plate with dimensions of 60 mm × 50 mm × 5 mm.

6. A method for in-situ manufacturing of high-performance Ti2AlNb alloy by laser directed energy deposition through particle size control according to claim 1, characterized in that In step S4, in the LDED in-situ manufacturing, the laser source uses a wavelength of 1064 nm, a spot diameter of 3.0 mm, and the laser system is IPGYLS-10000.

7. A method for in-situ fabricating a high-performance Ti2AlNb alloy by laser directed energy deposition with particle size control according to claim 1, characterized in that, In step S4, the powder feeding rate is 5.6 g / min.

8. A method for in-situ fabricating a high-performance Ti2AlNb alloy by laser directed energy deposition with particle size regulation according to claim 1, characterized in that, In step S4, the multi-scale simulation method includes a micro-scale laser scanning model and a macro-scale laser scanning model, and specifically includes the following steps: S41: Establish a micro-scale laser scanning model composed of the substrate TA15 and the first layer of powder; S42: In the micro-scale laser scanning model, adopt a graded grid division, use a fine grid near the deposited layer, and a coarse grid in other areas to obtain the molten pool size and temperature history of the single-track scanning trajectory points; S43: Select the middle position of the first layer of powder to obtain the molten pool size and temperature history, and read the time when the powder rises from the initial temperature to the highest temperature from the temperature history; S44: Simplify the laser heat source into an equivalent volumetric heat source with a length equal to the spot diameter, a width equal to the melt channel width, and a height equal to the molten pool depth, establish a macro-scale laser scanning model of a single layer of powder, calculate the equivalent volumetric heat source according to the powder heating time and molten pool size obtained from the micro-scale laser scanning model, and calculate the single-layer cooling time through the actual deposition time of a single layer, so as to control the thermal cycle of the entire component.

9. A method for in-situ fabricating a high-performance Ti2AlNb alloy by laser directed energy deposition through particle size control according to claim 8, characterized in that, In step 43, the time for the powder to rise from the initial temperature to the highest temperature is 0.1 s, and the molten pool size is a molten pool length of 0.6 mm, a melt channel width of 2.8 mm, and a molten pool depth of 0.7 mm.

10. A method for in-situ manufacturing a high-performance Ti2AlNb alloy by laser directed energy deposition with particle size regulation according to claim 9, characterized in that, The calculation formula for the equivalent volumetric heat source is: Wherein, A is the laser absorption rate, P is the laser power, ds is the laser spot diameter, dm is the width of the molten channel, and h is the depth of the molten pool; The single-layer cooling time is the single-layer actual deposition time minus the time for the powder to rise from the initial temperature to the highest temperature. The single-layer actual deposition time is 8 s, so the single-layer cooling time is 7.9 s.

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