CN117888005A - High-strength aluminum alloy suitable for laser powder bed melting manufacture - Google Patents
High-strength aluminum alloy suitable for laser powder bed melting manufacture Download PDFInfo
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- 229910000838 Al alloy Inorganic materials 0.000 title claims abstract description 145
- 239000000843 powder Substances 0.000 title claims abstract description 63
- 238000004519 manufacturing process Methods 0.000 title claims abstract description 23
- 238000002844 melting Methods 0.000 title claims abstract description 19
- 230000008018 melting Effects 0.000 title claims abstract description 19
- 239000013078 crystal Substances 0.000 claims abstract description 26
- 229910052782 aluminium Inorganic materials 0.000 claims abstract description 12
- XAGFODPZIPBFFR-UHFFFAOYSA-N aluminium Chemical compound [Al] XAGFODPZIPBFFR-UHFFFAOYSA-N 0.000 claims abstract description 9
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- 229910052726 zirconium Inorganic materials 0.000 abstract description 17
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- AZDRQVAHHNSJOQ-UHFFFAOYSA-N alumane Chemical class [AlH3] AZDRQVAHHNSJOQ-UHFFFAOYSA-N 0.000 description 1
- LUKDNTKUBVKBMZ-UHFFFAOYSA-N aluminum scandium Chemical compound [Al].[Sc] LUKDNTKUBVKBMZ-UHFFFAOYSA-N 0.000 description 1
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- Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
- Y02—TECHNOLOGIES OR APPLICATIONS FOR MITIGATION OR ADAPTATION AGAINST CLIMATE CHANGE
- Y02P—CLIMATE CHANGE MITIGATION TECHNOLOGIES IN THE PRODUCTION OR PROCESSING OF GOODS
- Y02P10/00—Technologies related to metal processing
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Abstract
The invention discloses a high-strength aluminum alloy suitable for laser powder bed melting manufacture, which comprises the following components in percentage by mass:mn 5.50-6.50%, mg 2.10-2.50%, zr0.50-0.90%, hf 0.90-1.60% and the balance of aluminum, wherein the Al-Mn-Mg-Zr-Hf alloy powder prepared by an air atomization method is formed by adopting a laser powder bed in a melting way. The aluminum alloy of the invention adopts Hf to completely replace expensive Sc, and combines the rapid solidification performance of the laser powder bed melting process to ensure that the Hf and Zr elements and Al formed by an Al matrix 3 The (Zr, hf) phase is used as a nucleating agent to promote the remarkable refinement of crystal grains in the aluminum alloy, form a crystal grain strengthening effect, improve the strength of the aluminum alloy, reduce the raw material cost for manufacturing the aluminum alloy by additive development, and promote the large-scale application of the aluminum alloy in the fields of aviation, aerospace and automobiles.
Description
Technical Field
The invention belongs to the technical field of metal materials, and particularly relates to a high-strength aluminum alloy suitable for laser powder bed melting manufacture.
Background
The laser powder bed fusion technology (L-PBF) is one of the most widely applied high-precision forming technologies in the field of metal additive manufacturing at present, firstly, a three-dimensional model of a part required by Computer (CAD) aided design is adopted and converted into a layered STL file, then layered slicing and scanning path planning of each layer are carried out, and then, laser beams with high energy density are adopted to melt the powder uniformly paved on a powder bed layer by layer, so that a complete part is formed.
Among the existing metallic materials for L-PBF, the use of aluminum alloys is limited. On the one hand, the near-eutectic Al-Si alloy with excellent L-PBF processing performance has the yield strength of 225 MPa-310 MPa in a deposition state and a heat treatment state, the ultimate tensile strength of 270 MPa-440 MPa and the elongation at break of 4% -10%, so that the mechanical performance of the Al-Si alloy is difficult to meet the use requirements of key force-bearing components in the fields of aerospace and the like. On the other hand, the direct use of aluminum alloys developed by conventional casting or forging processes for L-PBF suffers from a number of problems such as insufficient strength, susceptibility to cracking, and high porosity. Studies have shown that crack-free 2xxx and 7xxx series aluminum alloys can be produced using a combination of low laser power and low scan speed or powder bed preheating at 1500 ℃. However, it is difficult to obtain uniform cooling conditions when forming large-sized parts, especially in geometrically complex parts for engineering applications. In addition, the nucleating agent is directly introduced or formed in situ in the 2xxx or 7xxx aluminum alloy, and can be used as a nucleation site for alpha-Al crystallization due to low lattice mismatch degree of the nucleating agent and a matrix, so that the formation of superfine grains is facilitated, the strain generated in the L-PBF process is adapted, and the generation of solidification cracks is avoided. However, the mechanical properties obtained with this method are much lower than with corresponding wrought alloys. This is because the above alloys are not designed for L-PBF, the advantages of the unbalanced solidification characteristics of the L-PBF process are underutilized, and the corresponding heat treatment processes for these alloys are not fully applicable.
Air companies developed a custom for laser additive manufacturingThe alloy has the component composition of Al-4.5Mg-0.51Mn-0.66Sc-0.37Zr, and the modified aluminum alloy with hypereutectic aluminum scandium component is obtained by utilizing the rapid solidification characteristic of laser powder bed melting, the yield strength exceeds 500MPa, the tensile strength exceeds 520MPa, but Sc has the cost which is too high and is unfavorable for mass production of engineering, so researchers aim at L-PBF forming>Modification study of alloy.
Therefore, there is a need to develop a high strength aluminum alloy suitable for the L-PBF process and to customize the heat treatment process specifically for it to achieve good mechanical properties to meet the needs of engineering applications.
Disclosure of Invention
The technical problem to be solved by the invention is to provide a high-strength aluminum alloy suitable for laser powder bed melting manufacture aiming at the defects in the prior art. The aluminum alloy adopts Hf to completely replace expensive Sc, and combines the rapid solidification characteristic of a laser powder bed melting process to ensure that Zr, hf and Al matrix form Al 3 The (Zr, hf) phase is used as a nucleating agent to promote the remarkable refinement of crystal grains in the aluminum alloy and improve the strengthening effect of crystal boundaries, thereby improving the strength of the aluminum alloy and solving the problem of higher cost of the existing high-strength aluminum alloy.
In order to solve the technical problems, the invention adopts the following technical scheme: the high-strength aluminum alloy suitable for laser powder bed melting manufacture is characterized by comprising the following components in percentage by mass: mn 5.50-6.50%, mg 2.10-2.50%, zr 0.50-0.90%, hf0.90-1.60%, and the balance of aluminum, according to the design components of the target product aluminum alloy, preparing Al-Mn-Mg-Zr-Hf aluminum alloy powder by adopting a vacuum induction gas atomization method, and carrying out fusion forming by adopting a laser powder bed to obtain the aluminum alloy, wherein the room temperature tensile strength of the aluminum alloy in a deposited state is more than 520MPa, the yield strength is more than 460MPa, the elongation after break is more than 12%, the tensile strength in a heat treatment state is more than 570MPa, the yield strength is more than 520MPa, and the elongation after break is more than 5%.
The Al-Mn-Mg-Zr-Hf alloy designed and developed by the invention adopts Hf to completely replace Sc, and because Hf has larger solid solubility under the condition of rapid solidification, breaks through the solid solubility limit (0.18%) under the balance condition, has lower balance solid solubility and slower diffusion rate at the aging temperature of 300 ℃ which is commonly used in aluminum alloy, the Hf and the Al matrix form Al 3 Hf phase, al 3 Hf phase and other phases such as Al 3 The Zr phase has smaller lattice mismatch degree, better thermal stability and higher shear modulus, so that the aluminum alloy can maintain high strength performance at higher temperature; simultaneously increasing Zr content to enable Zr, hf and Al matrix to form Al 3 (Zr, hf) phase, a large amount of primary Al 3 The (Zr, hf) promotes the remarkable refinement of crystal grains in the aluminum alloy by increasing nucleation sites, and forms high-density crystal boundaries to forcefully block dislocation movement, so that the strength of the aluminum alloy is effectively improved, the coarsening resistance of the aluminum alloy is improved, the lattice mismatch with Al is reduced, the coarsening effect of sediment is weakened, the strengthening effect of a second phase is enhanced, and the strength of the aluminum alloy is further improved.
Compared withThe alloy of the invention is added with more Mn, mn is used as a solid solution reinforcing agent after quick solidification and a precipitation reinforcing agent after heat treatment, and plays a role in solid solution reinforcement in the solidification process and after heat treatment of the preparation of the aluminum alloy, thereby improvingThe strength of the aluminum alloy is improved, and meanwhile, the Al formed by other elements and an Al matrix is improved 3 Thermal stability of the X precipitate phase. In addition, a large amount of solute atoms Mg added into the aluminum alloy is also dissolved in the Al matrix in a solid solution manner, thereby playing a solid solution strengthening role, causing the interaction between a local strain field and dislocation, preventing dislocation from moving freely and further improving the strength of the aluminum alloy.
In addition, the elastic modulus of the alloy element added in the aluminum alloy is higher than that of the Al matrix, so that the elastic modulus of the aluminum alloy is improved, and the aluminum alloy has better rigidity. The elastic modulus of the high-strength aluminum alloy in the deposited state exceeds 75GPa, and the elastic modulus of the heat-treated state exceeds 80GPa.
The high-strength aluminum alloy suitable for laser powder bed melting manufacture is characterized by comprising the following components in percentage by mass: mn 6.14%, mg 2.12%, zr 0.90%, hf 1.15%, the balance being aluminum.
The high-strength aluminum alloy suitable for laser powder bed melting manufacture is characterized in that the grain size of the Al-Mn-Mg-Zr-Hf alloy powder is 20-70 mu m. The Al-Mn-Mg-Zr-Hf alloy powder with concentrated particle size distribution and consistent particle size reduces the phenomena of powder spheroidization and agglomeration in the laser powder bed melting preparation process, has higher surface finish, and ensures the consistency and uniformity of the laser powder bed melting process.
The high-strength aluminum alloy suitable for laser powder bed melting manufacture is characterized in that the forming parameters of the laser powder bed melting process are as follows: the laser power is 350W, the laser scanning speed is 1200mm/s, the scanning interval is 120 mu m, the powder layer thickness is 30 mu m, and the interlayer rotation angle is 67 degrees.
The high-strength aluminum alloy suitable for laser powder bed melting manufacture is characterized in that an aging heat treatment system adopted by the aluminum alloy in a heat treatment state is as follows: preserving heat for 4h at 375 ℃, and air cooling. The invention ensures that a large amount of nano/submicron rod-shaped Al is formed in the aluminum alloy by controlling the aging heat treatment system of the aluminum alloy 6 Mn precipitate phase and Al with average diameter of 2.5nm + -0.4 nm 3 (Zr, hf) nano-precipitants, thereby improving aluminumPeak hardness of alloy as deposited; meanwhile, the aging heat treatment system is only required to be carried out at a medium temperature, and is simpler and easier to realize.
The high-strength aluminum alloy suitable for laser powder bed melting manufacture is characterized by having a multi-stage heterogeneous structure: the central structure of the molten pool is micron-sized coarse columnar crystal CCG with the average grain size of 2.83 mu m plus or minus 1.57 mu m, and medium equiaxed crystal MEG with the average grain size of 1.94 mu m plus or minus 1.36 mu m, and the boundary structure of the molten pool is mainly submicron-sized ultra-fine equiaxed crystal UFG with the average grain size of 0.56 mu m plus or minus 0.01 mu m. In the process of the laser powder bed melting preparation of the aluminum alloy, superfine isometric crystals UFG (Ultra Fine Grain) are easy to form due to the larger cooling rate of the boundary of a molten pool, coarse columnar crystals CCG (Coarse Columnar Grain) are easy to form due to the smaller cooling rate of the inner part of the molten pool, and medium isometric crystals MEG (Medium Columnar Grain) between the two are also easy to form a multi-stage heterogeneous structure, wherein Zr, hf and an Al matrix added into the aluminum alloy form submicron primary Al 3 (Zr, hf) phase and distributed in the center of the superfine equiaxed crystal region, al with L12 structure 3 The lattice parameter of (Zr, hf) is 0.4126nm plus or minus 0.0104nm, the mismatch degree is low, and the lattice parameter has good coherent relation with the alpha-Al matrix, so that non-uniform nucleation sites are provided for the alpha-Al matrix, the grain refinement is promoted, and the grain boundary strengthening effect is improved; meanwhile, mn phase is uniformly distributed in the aluminum alloy in irregular block shape and long strip-shaped particles, and is precipitated along a grain boundary to play a role in precipitation strengthening, so that the strength of the aluminum alloy is improved.
The preparation process of the high-strength aluminum alloy comprises the following steps:
(1) Preparing aluminum alloy powder by adopting a vacuum induction gas atomization method according to the design components of the aluminum alloy of the target product, and detecting the components of the aluminum alloy powder by adopting an inductive coupling plasma optical emission spectrometer;
(2) Aluminum alloy powder with certain grain size is selected to be placed in a vacuum drying oven for drying treatment at 120 ℃ for 3 hours, then aluminum alloy modeling and layering slicing are carried out by adopting Magic software, and then laser powder bed melting technology is carried out by adopting German SLM-280HL equipment, wherein the equipment adopts double laser configuration, the laser wavelength of the prepared fiber laser is 1070nm, the laser spot diameter is 80 mu m, before forming, the Al6061 substrate is subjected to sand blasting treatment to reduce the surface roughness, then the powder is paved and melted and formed after heating to 150 ℃, the forming parameters are laser power 350W, the laser scanning speed is 1200mm/s, the scanning interval is 120 mu m, the thickness of the powder layer is 30 mu m, the interlayer rotation angle is 67 degrees, high-purity argon is adopted as protective gas in the forming process, and the forming process is completed in a glove box with the oxygen content of less than 600 ppm.
Compared with the prior art, the invention has the following advantages:
1. the Al-Mn-Mg-Zr-Hf aluminum alloy adopts Hf to completely replace expensive Sc, and combines the rapid solidification performance of a laser powder bed melting process to ensure that Zr, hf and Al formed by an Al matrix 3 The (Zr, hf) phase is used as a nucleating agent, the obvious refinement of crystal grains in the aluminum alloy is promoted by adding nucleation sites and heterogeneous nucleation, and the grain boundary strengthening and precipitation strengthening effects are improved, so that the strength of the aluminum alloy is improved, the raw material cost of the aluminum alloy is greatly reduced, the problem of higher cost of the existing aluminum alloy is solved, and the special aluminum alloy with high strength and controllable cost is provided for the melting and manufacturing of a laser powder bed, and is suitable for the aerospace field.
2. According to the aluminum alloy disclosed by the invention, a large amount of solute elements Mn with high equilibrium solubility and high expansion solubility under the condition of rapid solidification are added into an Al matrix to serve as a solid solution reinforcing agent after rapid solidification and a precipitation reinforcing agent after heat treatment, so that the reinforcing effect is effectively exerted, and the strength of the aluminum alloy is improved.
3. The invention prepares the nearly fully compact aluminum alloy forming piece through a laser powder bed melting process (L-PBF), and the aluminum alloy has excellent processing performance, no solidification cracks or obvious metallurgical defects, and is suitable for preparing parts with complex structures.
4. The invention adopts a medium-temperature aging heat treatment system which is simpler than other additive manufacturing aluminum alloys, so that alloy elements with lower equilibrium solid solubility in Al added in the aluminum alloys generate tiny dispersed sediment after heat treatment, the sediment strengthening effect is exerted, and the strength of the aluminum alloys is improved.
5. The aluminum alloy prepared by the laser powder bed melting technology has the multi-stage heterogeneous structure of micron-sized coarse columnar crystals CCG, medium equiaxed crystal MEG and submicron-sized superfine equiaxed crystal UFG, and primary Al distributed in the center of the superfine equiaxed crystal zone 3 The (Zr, hf) phase provides non-uniform nucleation sites for the Al matrix, promotes grain refinement, and improves grain boundary strengthening, thereby improving the strength of the aluminum alloy.
The technical scheme of the invention is further described in detail through the drawings and the embodiments.
Drawings
FIG. 1a is a graph (1000X) showing the morphology of the aluminum alloy powder prepared in example 1 of the present invention.
FIG. 1b is a graph (10000X) showing the morphology of the aluminum alloy powder prepared in example 1 of the present invention.
FIG. 2 is a microstructure of an aluminum alloy as deposited prepared in example 1 of the present invention.
FIG. 3a is an EBSD of the as-deposited aluminum alloy prepared in example 1 of the present invention.
FIG. 3b is an EBSD chart of the heat treated state of the aluminum alloy prepared in example 1 of the present invention.
FIG. 4a is a graph showing the engineering stress-strain curve of the aluminum alloy prepared in example 1 of the present invention.
FIG. 4b is a bar graph of tensile properties of the aluminum alloy prepared in example 1 of the present invention.
FIG. 5 is a STEM-EDS diagram of an aluminum alloy as deposited prepared in example 1 of the present invention.
Fig. 6a is a view of HAADF of an ultra-fine grain region in a as-deposited aluminum alloy prepared in example 1 of the present invention and a corresponding EDS plot.
Fig. 6b is an enlarged view of a portion of the broken line portion of fig. 6 a.
Fig. 6c is a HRTEM image and a corresponding FFT image of the as-deposited aluminum alloy prepared in example 1 of the present invention.
Detailed Description
Example 1
The high-strength aluminum alloy of the embodiment comprises the following components in percentage by mass: mn 6.14%, mg 2.12%, zr 0.90%, hf 1.15%, the balance being aluminum.
The preparation method of the high-strength aluminum alloy comprises the following steps:
step one, preparing aluminum alloy powder by adopting a vacuum induction gas atomization method according to design components of a target product aluminum alloy, and detecting the components of the aluminum alloy powder by adopting an inductive coupling plasma optical emission spectrometer, wherein the result is as follows: mn 6.10%, mg 2.22%, zr 0.81%, hf 1.09%, the balance being aluminum; the particle size detection of the aluminum alloy powder is carried out according to GB/T-19077 particle size distribution laser diffraction method, and the result is as follows: d10 As shown in fig. 1a and 1b, the aluminum alloy powder is mostly spherical with only a small amount of satellite powder;
(2) Drying aluminum alloy powder with the grain diameter of 20-70 μm in a vacuum drying oven at 120 ℃ for 3 hours to remove moisture in the aluminum alloy powder, modeling and slicing the aluminum alloy by adopting Magic software, and then forming a test block by adopting a German SLM-280HL device through a laser powder bed melting process.
Through detection, the density of the deposited aluminum alloy prepared in the embodiment is more than 99.7%, an SEM sample is prepared by adopting a standard metallographic sample preparation method, and the SEM sample is placed under a scanning electron microscope for microscopic structure observation, and the result is shown in figure 2.
FIG. 2 is a microstructure of an aluminum alloy in a deposited state prepared in accordance with the present embodiment, as can be seen from FIG. 2, the aluminum alloy has a plurality of stages of non-uniformity in the deposited stateStructure of: the central organization of the molten pool is micron-sized coarse columnar crystal CCG (b) 1 At), mesomeric MEG (b) 2 Where) the molten pool Boundary (BD) structure is mainly submicron ultra-fine equiaxed grain UFG (b) 3 At) a location.
The deposited aluminum alloy prepared in this example was mechanically polished and then subjected to stress relief treatment using a Leica ion polisher, and subjected to EBSD test under a focused ion/electron dual beam system, the results of which are shown in fig. 3 a.
Fig. 3a is an EBSD diagram of an aluminum alloy deposited state prepared in this embodiment, and it can be seen from fig. 3a that both micron-sized coarse columnar crystals CCG and medium equiaxed crystals MEG are located inside a molten pool, and submicron-sized ultra-fine equiaxed crystals UFG are located at the bottom of the molten pool; the area fraction of the CCG area and the MEG area in the molten pool is about 43 percent based on the EBSD chart, and the area fraction of the UFG area at the bottom of the molten pool is about 57 percent.
The deposited aluminum alloy prepared in the embodiment is insulated for 4 hours at 375 ℃, and air-cooled for aging heat treatment to obtain heat-treated aluminum alloy; meanwhile, the as-deposited aluminum alloy prepared in the embodiment is respectively subjected to heat preservation at 375 ℃ for 3 hours, 8 hours and 10 hours to serve as a comparison heat treatment state aluminum alloy. According to GB/T228-2002 standard of a room temperature tensile test method for metallic materials, the deposited aluminum alloy and the heat-treated aluminum alloy prepared in the embodiment and the control heat-treated aluminum alloy are processed and prepared into samples respectively, and then room temperature tensile test is carried out on an INSTRON 3382 electronic universal material tester, the displacement rate is 1mm/min, the results are shown in FIG. 4a and FIG. 4b, and the results show that: the aluminum alloy subjected to heat treatment at 375 ℃ for different time has high stress-strain curve coincidence ratio, the yield strength and the tensile strength are close, and the elongation is slightly different, wherein the aluminum alloy in a heat treatment state obtained by heat treatment at 375 ℃ for 4h has the best comprehensive mechanical property, and compared with the aluminum alloy in a deposition state, the aluminum alloy in a room temperature tensile strength of 520MPa, the yield strength of 464MPa, the elongation after break of 12%, the elastic modulus of 78GPa, the aluminum alloy in a heat treatment state has the tensile strength of 578MPa, the yield strength of 525MPa, the elongation after break of 5% and the elastic modulus of 83GPa.
Fig. 3b is an EBSD plot of the heat treated aluminum alloy prepared in this example, based on which the area fraction of the CCG and UFG regions inside the molten pool was counted to be about 43%, and the area fraction of the UFG region at the bottom of the molten pool was counted to be about 57%.
Comparing fig. 3a with fig. 3b shows that the grain size of the aluminum alloy in the heat treated state is not significantly changed compared with that in the deposited state, which means that the grains of the aluminum alloy are not grown under the heat treatment condition.
FIG. 5 is a STEM-EDS diagram of the aluminum alloy prepared in this example, and it can be seen from FIG. 5 that Mn and Mg phases are uniformly distributed in the aluminum alloy, mn phases precipitate particles and strip-shaped precipitates along grain boundaries, and Mg phases also segregate at the grain boundaries; the Zr and Hf elements are obviously separated out only in the superfine crystal area, and the separation size is smaller.
FIG. 6a is a HAADF image and a corresponding EDS image of an ultrafine grain region in an as-deposited aluminum alloy prepared in this example, FIG. 6b is a partial enlarged view of a dotted line portion in FIG. 6a, FIG. 6c is a HRTEM image and a corresponding FFT image of an as-deposited aluminum alloy prepared in this example, and it is understood from the combination of FIGS. 6a to 6c that a diamond-like precipitate phase enriched in Zr and Hf is observed in the center of grains in the ultrafine grain region, the average size thereof is about 87 nm.+ -. 29nm, and it can be determined from the diffraction spots that Al 3 (Zr, hf) phase.
Example 2
The high-strength aluminum alloy of the embodiment comprises the following components in percentage by mass: mn 6.46%, mg 2.35%, zr 0.60%, hf 1.58%, the balance being aluminum.
The preparation method of the high-strength aluminum alloy of the present embodiment is different from that of embodiment 1 in that the laser powder bed melting process parameters are as follows: the laser power is 325W, the laser scanning speed is 1000mm/s, and finally, the test block and the substrate are cut by adopting linear cutting to obtain square aluminum alloy with the dimensions of 10mm multiplied by 10mm, namely the deposited aluminum alloy; and then taking the deposited aluminum alloy, preserving heat for 4 hours at 375 ℃, and carrying out air cooling for aging heat treatment to obtain the heat-treated aluminum alloy.
According to the standard of GB/T228-2002 'room temperature tensile test method for metallic materials', a deposition state sample and a heat treatment state sample are respectively prepared by adopting the deposition state and heat treatment state aluminum alloy after aging heat treatment prepared in the embodiment, and then room temperature tensile test is carried out on an INSTRON 3382 electronic universal material tester, the displacement rate is 1mm/min, and the result shows that: the aluminum alloy has a room temperature tensile strength of 520MPa, a yield strength of 472MPa, a breaking elongation of 12%, an elastic modulus of 79GPa, a tensile strength of 584MPa in a heat treatment state, a yield strength of 525MPa, a breaking elongation of 5% and an elastic modulus of 82GPa.
Example 3
The high-strength aluminum alloy of the embodiment comprises the following components in percentage by mass: mn 5.56%, mg 2.48%, zr 0.51%, hf 0.94% and the balance being aluminum.
The preparation method of the high-strength aluminum alloy of the present embodiment is different from that of embodiment 1 in that the laser powder bed melting process parameters are as follows: the laser power is 325W, the laser scanning speed is 1200mm/s, and finally, the test block and the substrate are cut by adopting linear cutting to obtain square aluminum alloy with the dimensions of 10mm multiplied by 10mm, namely the deposited aluminum alloy; and then taking the deposited aluminum alloy, preserving heat for 4 hours at 375 ℃, and carrying out air cooling for aging heat treatment to obtain the heat-treated aluminum alloy.
According to the standard of GB/T228-2002 'room temperature tensile test method for metallic materials', a deposition state sample and a heat treatment state sample are respectively prepared by adopting the deposition state and heat treatment state aluminum alloy after aging heat treatment prepared in the embodiment, and then room temperature tensile test is carried out on an INSTRON 3382 electronic universal material tester, the displacement rate is 1mm/min, and the result shows that: the room temperature tensile strength of the aluminum alloy in a deposition state is 524MPa, the yield strength is 475MPa, the elongation after break is 12%, the elastic modulus is 78GPa, the tensile strength in a heat treatment state is 581MPa, the yield strength is 520MPa, the elongation after break is 5%, and the elastic modulus is 81GPa.
The above description is only of the preferred embodiments of the present invention, and is not intended to limit the present invention. Any simple modification, variation and equivalent variation of the above embodiments according to the technical substance of the invention still fall within the scope of the technical solution of the invention.
Claims (6)
1. The high-strength aluminum alloy suitable for laser powder bed melting manufacture is characterized by comprising the following components in percentage by mass: mn 5.50-6.50%, mg 2.10-2.50%, zr 0.50-0.90%, hf 0.90-1.60%, and the balance of aluminum, according to the design components of the target product aluminum alloy, preparing Al-Mn-Mg-Zr-Hf aluminum alloy powder by adopting a vacuum induction gas atomization method, and carrying out fusion forming by adopting a laser powder bed to obtain the aluminum alloy, wherein the room temperature tensile strength of the aluminum alloy in a deposited state is more than 520MPa, the yield strength is more than 460MPa, the elongation after break is more than 12%, the tensile strength in a heat treatment state is more than 570MPa, the yield strength is more than 520MPa, and the elongation after break is more than 5%.
2. The high-strength aluminum alloy suitable for laser powder bed fusion manufacturing according to claim 1, wherein the aluminum alloy comprises the following components in percentage by mass: mn 6.14%, mg 2.12%, zr 0.90%, hf 1.15%, the balance being aluminum.
3. The high-strength aluminum alloy suitable for laser powder bed fusion manufacture according to claim 1, wherein the grain size of the Al-Mn-Mg-Zr-Hf alloy powder is 20 μm to 70 μm.
4. A high strength aluminum alloy suitable for laser powder bed fusion manufacture according to claim 1, wherein the laser powder bed fusion process has forming parameters of: the laser power is 350W, the laser scanning speed is 1200mm/s, the scanning interval is 120 mu m, the powder layer thickness is 30 mu m, and the interlayer rotation angle is 67 degrees.
5. The high-strength aluminum alloy suitable for laser powder bed fusion manufacturing according to claim 1, wherein the aluminum alloy adopts an aging heat treatment system as follows: preserving heat for 4h at 375 ℃, and air cooling.
6. A high strength aluminum alloy suitable for laser powder bed fusion manufacture as claimed in claim 1, wherein the aluminum alloy has a multi-stage heterogeneous structure: the central structure of the molten pool is micron-sized coarse columnar crystal CCG with the average grain size of 2.83 mu m plus or minus 1.57 mu m, and medium equiaxed crystal MEG with the average grain size of 1.94 mu m plus or minus 1.36 mu m, and the boundary structure of the molten pool is mainly submicron-sized ultra-fine equiaxed crystal UFG with the average grain size of 0.56 mu m plus or minus 0.01 mu m.
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