A method for in-situ manufacturing high-performance Ti2AlNb alloy by particle size regulation and laser directed energy deposition

Abstract

This invention relates to the field of Ti2AlNb alloy manufacturing, and discloses a method for in-situ manufacturing of high-performance Ti2AlNb alloys through particle size-controlled laser-directed energy deposition (LDED). The method involves controlling the particle size of Ti, Al, and Nb powders, pre-mixing them using a planetary ball mill, and then using LDED to manufacture high-performance Ti2AlNb alloys in situ. This method allows for thorough melting and mixing of the three types of particles, thereby reducing element enrichment and depletion. Simultaneously, the phase compositions of the top and middle regions of the component are B2 and B2+α2+O phases, respectively, and the uneven heat conduction accelerates the cooling rate. The higher melting point Nb particles act as a heterogeneous nucleation substrate, and the heat released by the negative enthalpy of mixing alters the local heat flow direction before the solid / liquid interface, thereby changing the original grain growth direction and ultimately refining the microstructure. The average tensile strength at room temperature is 994.5 ± 46.29 MPa. Compared with components manufactured using pre-alloyed powder, the average UTS is increased by 14%, the grain refinement is 14%, the geometric dislocation density is increased by 60.8%, and the cost is reduced by approximately 80%.

Description

TECHNICAL FIELD

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

[0002] Ti2AlNb alloy is considered as a potential alternative material for nickel-based high-temperature alloy (8-8.5g / cm 3 ) in aerospace field due to its excellent high-temperature performance and moderate density (5-5.7g / cm 3 ). It can effectively reduce the weight of the structure and work stably at high temperature of 600-700℃, and has a wide application prospect. The manufacturing of Ti2AlNb alloy mainly depends on casting, hot mechanical processing (such as forging and rolling) and powder metallurgy technology, but these traditional methods have obvious limitations: the casting process is easy to cause coarse grains, Al burning and Nb segregation due to the differences in density and melting point of Ti, Al and Nb elements, thereby affecting the mechanical properties; the difficult deformation characteristics and high-temperature brittleness of the alloy limit its application of hot mechanical processing, although the development of powder metallurgy technology brings new opportunities for the manufacturing of Ti2AlNb-based alloy, but the expensive mold and high porosity hinder its application.

[0003] Additive manufacturing has the advantages of customized forming, low cost, high material utilization and short production cycle, and has a wide application scenario in Ti2AlNb alloy production. Because the plasticity of Ti2AlNb alloy at room temperature is low, it is difficult to make wire, so scholars mostly use multi-wire feeding process to arc additive manufacturing of Ti2AlNb alloy. Li et al. used TiNb-Al double-wire arc additive manufacturing of Ti2AlNb-based alloy, due to the composition and microstructure segregation, the UTS of the prepared alloy is only 504MPa, 0.41. Yu et al. improved the uniformity of arc additive manufacturing of Ti2AlNb alloy by optimizing the droplet transfer mode, the highest UTS is 891MPa, but coarse columnar crystals are prone to appear in the alloy. Polozov et al. used Ti, Al, Nb elemental powder as raw material, combined with selective laser melting (SLM) and hot isostatic pressing (HIP) and subsequent heat treatment, successfully prepared high-density Ti2AlNb alloy, but there are still un-melted Nb particles in the alloy. In another study, Polozov et al. annealed the Ti2AlNb alloy prepared in situ by SLM at high temperature, which improved the problem of un-melted niobium particles, but the UTS of the alloy was only 286MPa. Zhou et al. used pre-alloyed powder as raw material, and prepared high-performance Ti2AlNb alloy by SLM technology, the maximum UTS was 1090MPa. Polozov et al. used alloyed powder to manufacture Ti2AlNb alloy by SLM, and the prepared alloy had cracks, and the UTS was 693MPa. Zhou et al. prepared Ti2AlNb alloy by SLM process, and the UTS of the alloy after solid solution + aging treatment was 973MPa. Li et al. innovatively used double-wire parallel feeding technology to successfully prepare Ti2AlNb alloy by electron beam free forming, and the UTS was close to 879.5MPa. Using alloyed welding wire or powder can improve the performance of the alloy, but it will significantly increase the cost. Using single-element welding wire or powder to additively manufacture Ti2AlNb alloy is difficult to obtain uniform element distribution and good mechanical properties. Therefore, a powder particle size control model is proposed to manufacture Ti2AlNb alloy with excellent performance. SUMMARY

[0004] In view of the defects in the prior art, the purpose of the present application is to provide a method for in-situ manufacturing of high-performance Ti2AlNb alloy by particle size control laser directed energy deposition.

[0005] In order to achieve the above purpose, the present application provides the following technical scheme:

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

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

[0008] S2: after the screened Ti, Al, Nb powders are subjected to powder particle size regulation, they are pre-mixed by a planetary ball mill, and after being uniformly mixed, they are placed in a vacuum drying box at 120°C for drying for 120 minutes;

[0009] S3: polishing the surface of the substrate using sandpaper to remove the oxide layer, and then cleaning and drying using alcohol and acetone to remove residual impurities;

[0010] S4: placing the dried substrate under the LDED in-situ manufacturing device, introducing Ti, Al, and Nb mixed powders into the LDED in-situ manufacturing device through H2, and depositing layer by layer at a laser power of 1500W and a scanning rate of 5mm / s, and applying equivalent volume heat flux and cooling time to each deposition layer through multi-scale simulation to manufacture high-performance Ti2AlNb alloy.

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

[0012] Preferably, in step S2, the powder particle size regulation method is as follows: taking Nb powder as the standard, the diameter of Ti powder is 25-57μm, the average diameter is 41μm, the diameter of Al powder is 18-42μm, and the average diameter is 30μm.

[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 100r / min, and the mixing time is 180 minutes.

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

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

[0016] Preferably, in step S4, the powder feeding rate is 5.6g / 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: establishing 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, a gradual grid division is adopted, fine grid is used near the deposition layer and coarse grid is used in other areas, the molten pool size and temperature history of single scanning track point are obtained;

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

[0021] S44: The laser heat source is simplified as an equivalent volume heat source with a length of the spot diameter, a width of the melt width and a height of the molten pool depth, a macro-scale laser scanning model of single layer powder is established, the equivalent volume heat source is calculated according to the powder heating time and the molten pool size obtained by the micro-scale laser scanning model, and the single layer cooling time is calculated through the actual deposition time of single layer, so as to control the heat cycle of the whole component.

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

[0023] Preferably, the calculation formula of the equivalent volume heat source is:

[0024]

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

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

[0027] Compared with the prior art, the present application has the following beneficial effects:

[0028] In the present application, the Ti, Al and Nb powders are fully melted and mixed through the powder control model, so as to reduce the element enrichment and barren phenomenon; the heat cycle of each deposition layer is controlled, so as to realize that the phase composition of the top and middle regions of the component is B2 and B2+alpha2+O phase respectively, the non-uniform heat conduction accelerates the cooling rate, so as to increase the dislocation density; the Nb particles with high melting point are used as heterogeneous nucleation substrates, the heat released by the negative mixing enthalpy changes the local heat flow direction in front of the solid / liquid interface, so as to change the original grain growth direction, and finally the grains are refined.

[0029] The average tensile strength (UTS) of the in-situ manufactured Ti2AlNb alloy by LDED is 994.5 ± 46.29 MPa at room temperature, which is 14% higher than the pre-alloyed powder manufactured components, 14% finer grain size, 60.8% higher geometric dislocation density. The performance is comparable to the traditional process SPS and EBAM, and the cost is reduced by about 80%. BRIEF DESCRIPTION OF DRAWINGS

[0030] Figure 1 (a) The schematic diagram of the structure of the LDED device in the embodiment; (b) The schematic diagram of the longitudinal tensile test sampling; (c) The schematic diagram of the hardness test position; (d) The schematic diagram of the transverse tensile test sampling; (e) The schematic diagram of the organization characterization position.

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

[0032] Figure 3 The flow chart of the powder particle size regulation in the embodiment.

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

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

[0035] Figure 6 The microstructure and XRD analysis schematic diagram of the two powder preparation samples in the embodiment.

[0036] Figure 7 The TEM analysis schematic diagram of the two powder preparation samples in the embodiment.

[0037] Figure 8 The SEM-EDS element area scanning analysis schematic diagram of the two powder preparation samples in the embodiment.

[0038] Figure 9 The EBSD analysis schematic diagram of the two powder preparation samples in the embodiment.

[0039] Figure 10 The c and g area EBSD analysis schematic diagram of the two powder preparation samples in the embodiment.

[0040] Figure 11Schematic diagram of sample mechanical property analysis of two kinds of powders in the embodiment.

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

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

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

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

[0045] Figure 16 Cooling rate test result graph in the embodiment.

[0046] Figure 17 (a) is a schematic diagram of grain structure of Ti2AlNb alloy manufactured by pre-alloyed powder L-DED; (b) is a schematic diagram of Nb particle induced grain refinement; (c) is a schematic diagram of mixed enthalpy on grain growth.

[0047] Figure 18 Schematic diagram of tensile fracture morphology and overall fracture of the sample manufactured by two kinds of powders in the embodiment.

[0048] In the drawings: 1- planetary ball mill, 2- screen, 3- device body, 31- coaxial nozzle, 4- ABB mechanical arm, 5- laser device, 51- laser beam, 6- industrial computer, 7- powder feeder, 8- powder distributor, 9- work platform, 10- fabrication platform, 11- TA15 plate, 12- argon tank, 13- deposited layer. DETAILED DESCRIPTION

[0049] In order to make the objects, technical solutions and advantages of the present application clearer, the present application will be further described in detail below with reference to the drawings and embodiments. It should be understood that the specific embodiments described herein are only used to explain the present application and do not limit the present application. In addition, the technical features involved in each embodiment of the present application described below can be combined with each other as long as they do not conflict with each other.

[0050] Please refer to Figures 1-15 The present application provides a method for in-situ manufacturing high-performance Ti2AlNb alloy by particle size regulation and laser directed energy deposition, comprising the following steps:

[0051] S1: Screen Ti, Al, and Nb powders. The Ti, Al, and Nb powders should be spherical elemental powders, and the purity of Ti powder should be 99.8%, Al powder 99.9%, and Nb powder 99.9%.

[0052] S2: After adjusting the particle size of the selected Ti, Al, and Nb powders, they are pre-mixed in a planetary ball mill 1, and then dried in a vacuum drying oven at 120°C for 120 minutes.

[0053] Previous studies typically involved simply mixing powders of different particle sizes and using SLM to fabricate Ti2AlNb alloys in situ. However, this method often failed to produce Ti2AlNb alloys with uniform composition and poor mechanical properties. This powder particle size control method ensures that the mass ratio of individual Ti, Al, and Nb powder particles is close to the theoretical mass ratio of each element in the Ti2AlNb alloy (Ti:Al:Nb = 46:11:43). The powder particle size control process is as follows: Figure 3 As shown, when using a powder particle size control model ( Figure 3 (a) After determining the average particle size of Ti, Al, and Nb powder particles, select the corresponding sieve size 2( Figure 3 (b) Using Nb powder as a standard, Ti powder with a diameter of 25-57 μm and an average diameter of 41 μm was screened using molecular sieves, and Al powder with a diameter of 18-42 μm and an average diameter of 30 μm was screened. Finally, these were mixed in a planetary ball mill 1 at a speed of 100 r / min for 180 minutes. Figure 3 (c)).

[0054] Specifically, the mass of a single spherical powder particle 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 ρ represents the theoretical mass of Ti, Al, and Nb in the Ti2AlNb alloy. Ti ρ Al ρ Nb For the densities of Ti, Al, and Nb, d Ti d Al d Nb denoted as the average particle size of Ti, Al, and Nb elemental powders.

[0059] Since the Nb powder is expensive, the average particle size of Ti and Al is calculated to be 40.54 um and 29.84 um by formula (2) according to the particle size (average value is 31.99 um) of the purchased Nb powder. The screen mesh 2 aperture range is selected according to the average particle size of Ti and Al powder, which is 25-57 um (aperture average value is 41 um) and 18-42 um (aperture average value is 30 um) respectively. The final screened powder particle size is 41.52 um and 30.57 um respectively, which is close to the calculated value. The morphology and particle size distribution of Ti2AlNb pre-alloyed powder and Nb, Ti, Al elemental powder are as follows Figure 4 (a)-(h) as shown in, wherein Figure 4 (a, b, c, d) are SEM morphologies of powders Ti2AlNb (a), Nb (b), Ti (c), and Al (d); Figure 4 (e, f, g, h) are particle size distributions of Ti2AlNb (e), Nb (f), Ti (g), and Al (h); Figure 4 (i, j, k, l) are EDS results of mixed powder particles: mixed image (i), Ti (j), Al (k), and Nb (l). The results show that the Ti, Al, and Nb powder particles are uniformly distributed after mixing, and no obvious segregation or local missing phenomenon is observed.

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

[0061] S4: Place the dried substrate under the LDED in-situ manufacturing device. The laser source has a wavelength of 1064 nm and a spot diameter of 3.0 mm. The laser system is IPGYLS-10000. The Ti, Al, and Nb mixed powder is introduced into the LDED in-situ manufacturing device by H2. The powder feeding rate is 5.6 g / min. The laser power is 1500 W and the scanning rate is 5 mm / s. The equivalent body heat flux and cooling time are applied to each deposition layer to manufacture high-performance Ti2AlNb alloy. Finally, two 70-layer Ti2AlNb alloy thin-wall walls are obtained by additive manufacturing. The surface is observed to be metallic luster without cracks.

[0062] As Figure 1(a) shows that the LDED in-situ manufacturing device is located in the argon tank 12, and specifically includes the device body 3, the ABB mechanical arm 4, the laser device 5, the industrial computer 6, the powder feeder 7 and the powder distributor 8, the coaxial nozzle 31 is arranged on the device body 3, the ABB mechanical arm 4 controls the device body 3 to above the workbench 9, the manufacturing platform 10 is placed on the workbench 9, the TA15 plate 11 is placed on the manufacturing platform 10 in turn, the industrial computer 6 controls the laser device 5 to emit the laser beam 51 to the device body 3, and at the same time, the Ti, Al and Nb mixed powder in the powder feeder 7 is transported to the powder distributor 8 through Ar, and after being distributed, the powder is sprayed out of the coaxial nozzle 31 of the device body 3, and the deposition layer 13 of the Ti2AlNb alloy sample is formed on the TA15 plate 11.

[0063] The multi-scale simulation method in step S4 is used for efficient calculation of the LDED thermal cycle process of the Ti2AlNb alloy, and includes a micro-scale laser scanning model and a macro-scale laser scanning model, as shown in Figure 5 As shown in the micro-scale laser scanning model, the molten pool size and temperature history of the single scanning track point are obtained, so as to calculate the equivalent volume heat flux. In the macro-scale laser scanning model, the equivalent volume heat flux and cooling time are applied layer by layer, so as to complete the thermal cycle simulation of the whole component.

[0064] Specifically, the following steps are included:

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

[0066] S42: In the micro-scale laser scanning model, a gradual mesh division is adopted, a fine mesh is used near the deposition layer, and a coarse mesh is used in other areas, so as to obtain the molten pool size and temperature history of the single scanning track point;

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

[0068]

[0069] Wherein, P is the laser power, a f , a r respectively represent the ellipsoid half-axis length in the front and rear scanning direction of the laser beam, a f = 3mm, a r = 2mm, b and c respectively represent the half-axis length of the ellipsoid perpendicular to the scanning direction and the depth direction, b = 2.5mm and c = 0.4mm, f1 and f2 represent the energy distribution coefficients before and after the heat source f1 + f2 = 2, f1 = 0.4 and f2 = 1.6.

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

[0071] The laser parameters are constant during the deposition process, and the melt pool size fluctuates slightly, so the middle position of the first layer powder is selected to obtain the melt pool size and temperature history, and the results are as follows Figure 5 (b) and (c), under the action of laser, the powder undergoes instantaneous heating and cooling, and the time for the powder to rise from the initial temperature to the highest temperature is 0.1s, and the melt pool size is 0.6mm in length, 2.8mm in width and 0.7mm in depth.

[0072] S44: The laser heat source is simplified as an equivalent volume heat source with a length of the spot diameter, a width of the melt width and a height of the melt depth, a macro-scale laser scanning model of single layer powder is established, the equivalent volume heat source is calculated according to the powder heating time and the melt pool size obtained by the micro-scale laser scanning model, and the single layer cooling time is calculated through the actual deposition time of single layer, so as to control the thermal cycle of the whole 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 undergo similar heating and cooling processes as in the micro-scale laser scanning model, so 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 as an equivalent volume heat source with a length of the spot diameter, a width of the melt width and a height of the melt depth, and the calculation formula is as formula (4).

[0074] The calculation formula of the equivalent volume heat source is:

[0075]

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

[0077] The single layer cooling time is the actual deposition time of single layer minus the time for the powder to rise from the initial temperature to the highest temperature, and the equivalent volume heat source is calculated according to the powder heating time (0.1s) and the melt pool size obtained by the micro-scale laser scanning model, which is used to heat the single layer powder, as shown in Figure 5 (d), since the actual deposition time of single layer is 8 seconds, the single layer cooling time is set to 7.9s. Finally, the equivalent volume heat flux and cooling time are applied to each deposition layer layer by layer, so as to complete the thermal cycle control of the whole component, and the results are as follows Figure 5 (e).

[0078] The performance of the Ti2AlNb alloy manufactured in-situ by the LDED is verified by experimental analysis.

[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 longitudinal tensile properties (a) and transverse tensile properties (b). Figure 1 Figure 1 The tensile test was completed on a KPL-FocalU100 (China) device, and the test method referred to the Chinese standard GB / T228.1-2010 “Metallic materials-tensile testing at ambient temperature”, the strain rate was 1x10 -3 s -1 , the tensile rate was 0.5 mm / min, the strain was measured by a video extensometer, and the final UTS was the average value of three samples.

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

[0082] The samples taken from the middle and top of the Ti2AlNb alloy thin-walled wall were ground and polished, and then microstructure characterization was performed, as shown in Figure 1 (e). After the samples were corroded by Kroll reagent (HF:HNO3:H2O=1:6:10), the microstructure was observed by scanning electron microscope (SEM, ZEISS Gemini SEM 300, Germany), and the element distribution of the samples was analyzed by EDS. The phase composition was determined by X-ray diffraction (XRD, Rigaku Ultima IV, Japan). The EBSD characterization of the samples prepared by premixed powder and pre-alloyed powder was performed by EDA-TSL (USA) equipment, with an acceleration voltage of 15 kV, a beam current of 1 nA, and a working distance of 14 mm. The grain size and kernel average misorientation (KAM) analysis used a 1 μm step scan, and the O / α2 phase precipitation characteristic analysis used a 0.1 μm fine step scan. The EBSD data was analyzed by AZtecCrystal software, and the samples were prepared by vibration polishing. The phase of the microstructure was analyzed by 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 cycling XRD analysis of Ti2AlNb alloy samples

[0086] XRD analysis shows that region b of the deposited sample only contains B2 phase, while region c contains B2, O and a2 phases. In order to explore the phase evolution mechanism of the deposited sample, further analysis of the thermal cycle during deposition is needed. Select A point (located in region b) and B point (located in region c) in the finite element model for analysis, as shown in Figure 12 (a). Combined with the pseudo-binary phase diagram of Ti-22Al-xNb alloy (see Figure 12 (b)), when the temperature is higher than 1060℃, the phase composition of the alloy is single B2 phase; when the temperature is lower than 1060℃, a2 and O phases begin to precipitate, so 1060℃ is taken as the boundary, the lowest temperature of the thermal cycle is defined as high-temperature thermal cycle (HTTC) when it is higher than 1060℃; the highest temperature is defined as low-temperature thermal cycle (LTTC) when it is lower than 1060℃; if 1060℃ is between the lowest and highest temperatures, it is defined as medium-temperature thermal cycle (MTTC). It can be seen that A point has experienced 10 times of HTTC, B point has experienced 6 times of MTTC and 29 times of LTTC. With the progress of deposition, the temperature cumulative effect makes the peak temperature of B point (2477℃) to A point (2801℃) gradually increase, as shown in Figure 12 (c).

[0087] An infrared camera (FLIRA615) was used to measure the temperature field, and K-type thermocouples (KPS-IN600-K, 10Hz) were used to calibrate the infrared data before the experiment. The thermocouples were placed on the left side of the Ta15 substrate near the start of the deposition ( Figure 13 (a)). The calibration results show that the thermocouples and the infrared temperature measurement data are in good agreement ( Figure 13 (b)), which verifies the reliability of the infrared temperature measurement method. To verify the accuracy of the model, the surface temperature of the deposited sample was measured using an infrared camera, and the position of the selected point is as shown in Figure 13 (c), which corresponds to point A in the finite element model. The measured temperature was compared with the simulation results, and the results are shown in Figure 13 (d). A point has experienced HTTC, and the error between the simulated peak temperature (2658℃) and the measured temperature (2768℃) is 3.9%. Although there is a certain error between the simulation data and the test data, the temperature change trend of the simulation and the actual measurement is basically consistent.

[0088] IV. Microstructure and element distribution of pre-mixed powder and pre-alloyed powder deposited samples

[0089] (1) Microstructure and XRD analysis of pre-mixed powder and pre-alloyed powder

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

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

[0092] Based on this microstructure and XRD analysis, such as Figure 6 (b)-(d) and 6(f)-(h) respectively illustrate the microstructure of samples prepared from premixed powder and pre-alloyed powder, combined with XRD ( Figure 6 Analysis (i)-(k) revealed that the top regions (b, f) consisted only of the B2 phase, with clearly visible grain boundaries; while the middle (c, g) and bottom (d, h) regions were composed of B2, O, and α2 phases, exhibiting similar microstructures. Both regions showed O and α2 phases precipitated within the B2 matrix, displaying typical characteristics. Tissue characteristics. Samples obtained by both preparation methods showed the same tissue evolution pattern along the deposition height direction: from the bottom to the top, the tissue gradually transitioned from a three-phase structure of B2+O+α2 to a single B2 phase.

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

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

[0095] Figure 6 (a) shows the brightness field pattern of this region. Selected area electron diffraction (SAED) was performed on three characteristic regions (A, B, C) from this pattern, and the results are as follows: Figure 6 (b)-(c) correspond to the B2 matrix phase, O phase, and α2 phase, respectively. Both the O and α2 phases exhibit needle-like distributions, but the O phase is relatively smaller. The phase composition is consistent with the XRD test results. Furthermore, high-resolution images of the area near the grain boundaries of the α2 phase (region D) show further evidence of this. Figure 7 (e) Perform Inverse Fourier Transform (IFFT) (see...) Figure 7(f)) and a large number of dislocations and lattice distortions were found near the grain boundaries. The interaction and interlacing of dislocations in different directions made it difficult for them to move, achieving the effect of dislocation pinning, which ultimately effectively improved the tensile strength of Ti2AlNb alloy.

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

[0097] (3) SEM-EDS element area scanning analysis of pre-mixed powder and pre-alloyed powder

[0098] The element distribution around the melt pool of the pre-mixed powder sample was analyzed by SEM-EDS element area scanning analysis Figure 8 (a)-(d)), the results showed that the Nb element in the melt pool had completely melted, and no unmelted Nb particles were observed. Further high-magnification SEM-EDS analysis of the element distribution in the pre-mixed powder region c Figure 8 (e)-(h)) and the pre-alloyed powder region g Figure 8 (i)-(l)) of the prepared samples found that only a small amount of segregation occurred at the grain boundaries of the two samples, and there was no macroscopic segregation. The Ti, Al, and Nb elements of the pre-mixed powder prepared sample were uniformly distributed, the deviation of the actual mass ratio (47.4:10.3:42.3) from the theoretical mass ratio (46:11:43) was small, and no unmelted Nb particles were detected at the melt pool boundary, indicating that this method effectively solved the problem of uneven composition distribution of in-situ additive manufacturing of Ti2AlNb alloy using pre-mixed powder.

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

[0100] EBSD analysis results of the middle regions of the Ti2AlNb alloy samples prepared by pre-mixed powder and pre-alloyed powder Figure 9 regions c and g). The average grain size was 211.3 μm and 245.7 μm Figure 9 (b) and (f)), respectively. The grain size of the pre-mixed powder prepared sample was refined by 14%, and the grain size was less than 400 μm, while the grain size of the pre-alloyed powder prepared sample was greater than 400 μm Figure 9 (a)). The average orientation difference (Kernel Average Misorientation, KAM) of the two was also analyzed, and the results are shown in Figure 9(c) and (g), the KAM of the pre-mixed powder manufactured sample (1.39°) is significantly higher than that of the pre-alloyed powder sample (0.54°), the average misorientation is increased by 61.2% (d) and (h) are the pole figures of the pre-mixed powder manufactured sample and the pre-alloyed powder sample, respectively. The texture of the pre-mixed powder manufactured sample is relatively weakened compared with the pre-alloyed powder sample, and the grain preferred orientation is weakened. Figure 9 (e) and (i), the higher KAM value indicates that the dislocation density (GND) in the sample is higher. Figure 9 (i) and (k) are the pole figures of both, the texture of the pre-mixed powder manufactured sample is relatively weakened compared with the pre-alloyed powder sample, and the grain preferred orientation is weakened. B2 phase is the main plastic phase in Ti2AlNb alloy, which has more slip systems, among which {110}<111> and {112}<111> are the main slip systems. The Schmid factors of the pre-mixed powder and the pre-alloyed powder samples on these two slip systems are calculated: the results are as follows Figure 9 (j) and (l), the Schmid factor of the pre-mixed powder sample (0.45) is smaller than that of the pre-alloyed powder sample (0.47) on both {110}<111> and {112}<111> slip systems, indicating that the slip system activation of the pre-mixed powder sample is more difficult, which may be the reason for the improvement of tensile strength.

[0101] In addition, the microstructure, precipitate phase orientation and volume fraction of the pre-mixed powder (region c) and the pre-alloyed powder (region g) samples were analyzed by high-resolution EBSD (step size 0.1 μm). The results show that the microstructure of region c is thinner than that of region g Figure 10 (a), (d)). The precipitate phase of both samples shows 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 relationship of (001)O / / {0001}α2 and (001)O / / {110}B2, which is consistent with the classic Burgers orientation relationship. Figure 10 (c) and (f) are the phase composition diagrams of region c and region g, where B2 phase is green, O phase is purple, and α2 phase is yellow. According to the statistics of AZtecCrystal software, the contents of B2 phase, O phase and α2 phase of the pre-mixed (region c) powder sample are 72.7%, 9.1% and 18.2%, respectively; while the contents of B2 phase, O phase and α2 phase of the pre-alloyed (region f) powder sample are 71.6%, 25.2% and 3.1%, respectively, as shown in Figure 10(h) It is worth noting that, although both the a2 phase in premixed (region c) and pre-alloyed (region g) samples exhibit a needle-like morphology, there is a significant difference in the morphology of the O phase between the two: the O phase in the premixed sample is needle-like, while the O phase in the pre-alloyed sample is lamellar. A faster cooling rate will promote the thinning of the O phase lamellar structure, reduce the volume fraction of the O phase, and induce the precipitation of needle-like O phase in the matrix. Since the cooling rate of the premixed powder sample is higher than that of the pre-alloyed powder, the volume fraction of the O phase is lower, the microstructure is thinner, and the needle-like precipitated phase is formed in the B2 matrix, which is completely consistent with the EBSD analysis.

[0102] (5) Mechanical property analysis of premixed and pre-alloyed powders

[0103] The tensile test results show that the average UTS and average EL of the in-situ manufactured samples of the premixed powder are 973.75 ± 36.85 MPa, 1.07 ± 0.1% and 1016 ± 55.72 MPa, 1.11% ± 0.2% (a) and (c) respectively; while the average UTS and average EL of the samples manufactured from pre-alloyed powder are 851.24 ± 20.75 MPa, 0.85 ± 0.03% and 849 ± 43.2 MPa, 0.89% ± 0.1% (b) and (d) respectively. The results show that the average UTS and EL of the premixed powder samples are superior to those of the pre-alloyed powder samples. Figure 11 Figure 11

[0104] In addition, the hardness test results along the deposition direction (6 mm, 5 equidistant positions, 3 test points in the horizontal direction for each position) (e) show that the average hardness of the B2+O+ a2 phase region of the premixed powder and pre-alloyed powder samples is 436 HV and 424 HV respectively, and the average hardness of the B2 phase region is 363 HV and 361 HV respectively. The hardness distribution trend of the two powder prepared samples is basically the same, and the hardness of the B2+O+ a2 region is higher than that of the B2 phase region, which is due to the existence of a2 / O phase as a strengthening phase and B2 phase as a plastic phase in Ti2AlNb alloy. However, the hardness values in the horizontal and vertical directions in the same phase composition region are small, which confirms that the overall performance of the premixed powder sample is uniform. Figure 11

[0105] In order to evaluate the position of the LDED in-situ manufacturing process in commonly used processes, we compared its mechanical properties with those of components manufactured by commonly used processes, such as Figure 11 ​​​(f) shown. The results show that the UTS of the LDED in-situ manufactured component is close to SLM and SPS (Spark Plasma Sintering), and is superior to PM (Powder Metallurgy) and WAAM (Wire Arc Additive Manufacturing). At present, the process of in-situ manufacturing Ti2AlNb alloy with the best performance is EBAM (Electron Beam Additive Manufacturing), although Polozov et al. also use SLM to in-situ manufacture Ti2AlNb alloy, but the mechanical properties are poor. The mechanical properties of the components manufactured by the LDED in-situ manufacturing process proposed by the present application are close to EBAM, and the cost is lower. However, compared with the components manufactured by forging and part of the SLM process, the mechanical properties still have a gap, especially in plasticity.

[0106] Five, alloy element homogenization and phase transformation mechanism analysis

[0107] (1) Analysis of three kinds of powder homogenization

[0108] After the three kinds of powder are mixed uniformly and sent into the molten pool, the powder is randomly distributed. If the powder particle size is not optimized, it is easy to cause uneven distribution of elements. Taking Nb element as an example for analysis, as shown in Figure 14 (a), the Nb powder particle size of region 1 is small (small in mass), which 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 element in other regions to flow to region 1, thereby causing the generation of a Nb-poor region. The Nb powder particle size of region 2 is large (large in mass), which is easy to produce melted Nb powder. Even if it is fully melted, the remaining Nb element cannot flow to other regions in time, and finally forms a Nb-rich region, which is consistent with the observation of Polozov, Grigoriev and Wang et al.

[0109] The present application optimizes the particle size of Ti, Al, Nb powder particles, so that the mass ratio of Ti, Al, Nb single powder is close to the theoretical mass ratio of each element of Ti2AlNb alloy (Ti:Al:Nb=46:11:43), and the average particle size of the three kinds of powder is 41.52um, 30.57um and 31.99um respectively. The Nb powder particle size is moderate and is easy to melt. As shown in Figure 14 (b), when the Ti, Al, Nb three adjacent powders satisfy the particle size ratio, the in-situ generated Ti2AlNb alloy satisfies the theoretical mass ratio, and there is no lack or excess of Nb element, thereby reducing the probability of occurrence of element deficiency or enrichment.

[0110] (2) Analysis of phase transformation mechanism of three kinds of powder

[0111] As shown in Figure 15 When the laser acts on the pre-mixed powder, the powder melts rapidly, the temperature rises rapidly to the peak, and then falls rapidly with a rate of more than 103 The cooling rate of K / s decreases sharply, and the B2 phase cannot be precipitated in time, so that the Ti2AlNb alloy without thermal cycle experiences the phase composition of B2 phase, as shown in Figure 15 (a). From Figure 12 (c), it can be seen that point A is completely in the HTTC condition, and the O phase and the a2 phase cannot be precipitated, so that the phase composition is a single B2 phase. Point B has experienced 5 MTTC and 35 LTTC, and in the MTTC process, the a2 phase precipitates and dissolves alternately, and the dissolution temperature is higher than the precipitation temperature, which leads to the a2 phase not being obviously precipitated, and after entering the LTTC stage, when the thermal cycle temperature interval is 1010-1060℃, the a2 phase begins to precipitate along the grain boundary Figure 15 (b)), and when the temperature is lower than 1010℃, the strip-shaped O phase begins to precipitate Figure 15 (c)), as shown in Figure 13 (c), which finally leads to the phase composition of point B being B2, O and a2 phase. In addition, the O phase formation mechanism of the Ti2AlNb alloy component prepared by the premixed powder is described. In the present application, the precipitation sequence of the O phase is described as follows: first, the Ti, Al, Nb premixed particles are melted to form a liquid phase under the action of laser, and then the B2 phase is nucleated from the liquid phase during solidification. After LTTC, the O phase begins to precipitate, and an intermediate transition phase B19 will appear during the precipitation process, that is, B2→B19→O phase transition occurs.

[0112] The temperature curves of the first layer midpoint (point A1, A2) and the 35th layer midpoint (point B1, B2) of the premixed powder Figure 16 (a) and the pre-alloyed powder Figure 16 (b) are obtained by using an infrared camera, and the corresponding cooling rates are calculated, as shown in Figure 16 (c) and (d): the cooling rates of the first layer of the premixed powder and the pre-alloyed sample are 1563℃ / s and 898℃ / s, respectively; the cooling rates of the 35th layer are 934℃ / s and 795℃ / s, respectively. It can be found that the cooling rate of the premixed powder is always higher than that of the pre-alloyed powder, whether at the first layer or at the 35th layer midpoint. Faster cooling rate is beneficial to the formation of fine grain structure, which not only increases the number of grain boundaries, but also causes more significant lattice distortion. Since there is a difference in the orientation of the grains on both sides of the grain boundary, the dislocations generated by the slip in one side of the grain cannot be directly transmitted to the adjacent grain[3], which leads to the accumulation of dislocations and the significant increase of dislocation density, thereby triggering higher KAM value.

[0113] The grain morphology of the component prepared by the pre-alloyed powder is mainly columnar crystal and equiaxed crystal, as shown in Figure 17 (a); Figure 17(b) is a schematic diagram of heterogeneous nucleation. The Nb particles melt later than the Ti and Al particles under laser irradiation. Nb can act as an effective nucleation substrate for the B2 phase, promoting nucleation and refining the structure; the liquid mixing enthalpy of the Ti, Al, and Nb pre-mixed powder is less than zero, indicating that the alloying process is an exothermic reaction. The released heat will change the heat flow direction of the local area in front of the solid / liquid interface. When the next layer is deposited, the original grain growth along the deposition direction is interrupted, promoting the nucleation and growth of new grains in front of the solid / liquid interface, resulting in a refined structure, as shown in Figure 17 (c).

[0114] The pre-mixed powder sample has a higher KAM value than the pre-alloyed powder, indicating 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, thereby transforming 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 subsequent dynamic recrystallization, these low-angle grain boundaries will gradually evolve into high-angle grain boundaries (HAGB > 15°), ultimately promoting grain refinement.

[0115] The tensile fracture of the pre-mixed powder and pre-alloyed powder samples is shown in Figure 18 (a) and (b), where the pre-mixed powder sample exhibits a mixed mode characteristic: both intergranular fracture, transgranular fracture, and a small amount (about 7%) of quasi-cleavage fracture exist. It is worth noting that the small amount of cleavage fracture does not have enough toughening effect to offset the brittleness caused by coarse grains, so the overall still shows brittle fracture. The pre-alloyed powder sample exhibits typical brittle fracture characteristics, mainly showing intergranular fracture and transgranular fracture modes. In the intergranular fracture area Figure 18 (c) and (e)), the rock pattern structure formed by coarse equiaxed grains can be clearly observed, and there are obvious cracks between the grains. In the transgranular fracture area Figure 18 (a) and (f)), a large number of typical river patterns can be seen, with their direction perpendicular to the grain boundaries. The quasi-cleavage fracture morphology observed in the pre-mixed powder tensile fracture (as shown in Figure 18 (d)) is consistent with these characteristics, confirming that the fracture exhibits quasi-cleavage characteristics rather than a dimple structure. This fracture morphology is highly consistent with the results reported in the literature, both showing typical brittle fracture characteristics.

[0116] In summary: under the same conditions, compared with pre-alloyed powder, the advantages of using premixed powder LDED in-situ manufacturing Ti2AlNb alloy are mainly in the performance and cost of two aspects: compared with pre-alloyed powder manufacturing samples, the average UTS is increased by 12.5%, the average EL is increased by 20.6%, and the performance is equivalent to that of Ti2AlNb alloy manufactured by traditional process. Statistics show that the average price required to purchase 1kg Ti2AlNb pre-alloyed powder is $778.18, while the average price of 1kg premixed powder (Ti: 460g, Al, 110g, Nb: 430g) is $155.64, and the cost is reduced by 80%. It can be found that using premixed powder can manufacture high-performance Ti2AlNb alloy components at low cost, which is conducive to quickly promoting the commercial application of Ti2AlNb alloy.

[0117] The present application uses Ti, Al, Nb elemental powder as raw material to in-situ manufacture Ti2AlNb alloy with excellent performance by LDED:

[0118] (1) By powder control model, the mass ratio of Ti, Al, Nb single powder is close to the theoretical mass ratio of each element of Ti2AlNb alloy (Ti:Al:Nb=46:11:43), when the particle size ratio of Ti, Al, Nb three adjacent powders is satisfied, the in-situ generated Ti2AlNb alloy satisfies the theoretical mass ratio, and there is no lack or excess of elements, thereby reducing the probability of occurrence of element deficiency or enrichment. The mass of Ti, Al, Nb (47.4:10.3:42.3) in the deposited Ti2AlNb alloy component deviates less from the theoretical mass, the element distribution is uniform, and there is no obvious macrosegregation phenomenon.

[0119] (2) The top region of the component only experiences HTTC, and the phase composition is single B2 phase; the middle region is dominated by LTTC, and the phase composition is B2+α2+O phase.

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

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

[0122] The above merely describes the specific embodiments of the present application, but the technical features of the present application are not limited to this. Any changes or modifications made by those skilled in the art within the scope of the present application are covered by the patent range of the present application.

Claims

1. A method for in-situ manufacturing high-performance Ti2AlNb alloys by particle size-regulated laser directed energy deposition, characterized in that, Comprise the following steps: S1: screening Ti, Al, Nb powder, and ensuring that the purity of Ti powder is 99.8%, the purity of Al powder is 99.9%, and the purity of Nb powder is 99.9%; the Ti, Al, Nb powder is spherical single-phase powder, S2: after powder particle size regulation of the screened Ti, Al, Nb powder, pre-mixing is carried out through a planetary ball mill, the ratio of Ti, Al, Nb powder entering the planetary ball mill is 46:11:43, and after mixing, it is placed in a vacuum drying box at 120 DEG C for drying for 120 minutes; The powder particle size regulation method is that: taking Nb powder as a standard, the diameter of Ti powder is 25-57 microns, the average diameter is 41 microns, the diameter of Al powder is 18-42 microns, and the average diameter is 30 microns through molecular sieve screening; S3: polishing the surface of the substrate using sandpaper to remove the oxide layer, then cleaning and drying with alcohol and acetone to remove residual impurities; S4: placing the dried substrate under the LDED in-situ manufacturing device, introducing Ti, Al, Nb mixed powder into the LDED in-situ manufacturing device through H2, depositing layer by layer with a laser power of 1500 W and a scanning rate of 5 mm / s, and applying equivalent volume heat flux and cooling time to each deposition layer through multi-scale simulation to manufacture high-performance Ti2AlNb alloy; The multi-scale simulation method comprises a micro-scale laser scanning model and a macro-scale laser scanning model, and specifically comprises the following steps: S41: establishing a micro-scale laser scanning model composed of a substrate TA15 and a first layer of powder; S42: in the micro-scale laser scanning model, using a gradual grid division, using a fine grid near the deposition layer and a coarse grid in other areas, obtaining the molten pool size and temperature history of the single scanning track point; S43: selecting the middle position of the first layer of powder to obtain the molten pool size and temperature history, reading the powder heating time from the initial temperature to the highest temperature in the temperature history; S44: simplifying the laser heat source into an equivalent volume heat source with a length of the spot diameter, a width of the melt width, and a height of the molten pool depth, establishing a macro-scale laser scanning model of a single layer of powder, calculating the equivalent volume heat source according to the powder heating time and the molten pool size obtained from the micro-scale laser scanning model, and calculating the single layer cooling time through the actual deposition time of the single layer, so as to control the thermal cycle of the whole component.

2. The method of claim 1, wherein the method is characterized by, The rotation speed of the planetary ball mill in step S2 is 100 r / min, and the mixing time is 180 minutes.

3. The method of claim 1, wherein the method is characterized by, The substrate in step S3 is a TA15 plate with a size of 60 mm x 50 mm x 5 mm.

4. The method for in-situ manufacturing high-performance Ti2AlNb alloy by particle size controlled laser directed energy deposition according to claim 1, characterized in that, The laser source in the LDED in-situ manufacturing in step S4 adopts a wavelength of 1064 nm, a spot diameter of 3.0 mm, and an IPG YLS-10000 laser system.

5. The method for in-situ manufacturing high-performance Ti2AlNb alloy by particle size controlled laser directed energy deposition according to claim 1, characterized in that, The powder feeding rate in step S4 is 5.6 g / min.

6. The method for in-situ fabrication of high-performance Ti2AlNb alloy by particle size controlled laser directed energy deposition according to claim 1, characterized in that, The powder heating time 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 width of 2.8 mm, and a molten pool depth of 0.7 mm.

7. The method of claim 6, wherein the method is characterized by, The calculation formula of the equivalent volume heat source is: wherein, A is the laser absorption, P is the laser power, ds is the laser spot diameter, dm is the kerf width, h is the melt pool depth; 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 maximum temperature. The single layer actual deposition time is 8 s, so the single layer cooling time is 7.9 s.

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