A high-strength aluminum alloy suitable for laser powder bed melting manufacturing

Abstract

This invention discloses a high-strength aluminum alloy suitable for laser powder bed melting manufacturing, composed of the following mass percentages: Mn 5.50%–6.50%, Mg 2.10%–2.50%, Zr 0.50%–0.90%, Hf 0.90%–1.60%, with the balance being aluminum. This aluminum alloy is formed by laser powder bed melting of Al-Mn-Mg-Zr-Hf alloy powder prepared by gas atomization. In this invention, Hf completely replaces the expensive Sc. Combined with the rapid solidification performance of the laser powder bed melting process, the Al3(Zr,Hf) phase formed by Hf and Zr elements with the Al matrix acts as a nucleating agent, promoting significant grain refinement in the aluminum alloy, forming a grain strengthening effect, and improving the strength of the aluminum alloy. The development of this aluminum alloy reduces the raw material cost of additive manufacturing aluminum alloys and promotes the large-scale application of aluminum alloys in the aerospace and automotive fields.

Description

Technical Field

[0001] This invention belongs to the field of metal materials technology, and specifically relates to a high-strength aluminum alloy suitable for laser powder bed melting manufacturing. Background Technology

[0002] Laser powder bed melting (L-PBF) is one of the most widely used high-precision forming technologies in the field of metal additive manufacturing. First, computer-aided design (CAD) is used to design the three-dimensional model of the required part and convert it into a layerable STL file. Then, layer slicing and scanning path planning for each layer are performed. Finally, a high-energy-density laser beam is used to melt the powder uniformly laid on the powder bed layer by layer, thereby forming a complete part.

[0003] Among existing metallic materials used for L-PBF (Laminated Phthalochemical Bonding), the application of aluminum alloys is limited. On the one hand, near-eutectic Al-Si alloys, which possess excellent L-PBF processing properties, have yield strengths of 225 MPa–310 MPa in both deposited and heat-treated states, ultimate tensile strengths of 270 MPa–440 MPa, and elongation at break of 4%–10%. Therefore, the mechanical properties of Al-Si alloys are insufficient to meet the requirements of critical load-bearing components in aerospace and other fields. On the other hand, directly using aluminum alloys developed through traditional casting or forging processes for L-PBF faces many 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 manufactured by combining low laser power and low scanning speed or powder bed preheating at 1500°C. However, it is difficult to obtain uniform cooling conditions when forming large-sized parts, especially in geometrically complex parts used in engineering applications. Furthermore, directly introducing or in-situ forming nucleating agents into 2xxx or 7xxx aluminum alloys can serve as nucleation sites for α-Al crystallization due to the low lattice mismatch between the nucleating agent and the matrix. This promotes the formation of ultrafine grains, accommodating the strain generated during L-PBF and preventing solidification cracks. However, the mechanical properties obtained using this method are mostly lower than those of the corresponding forged alloys. This is because these alloys are not designed for L-PBF, the advantages of the non-equilibrium solidification characteristics of the L-PBF process are not fully utilized, and the corresponding heat treatment processes are not entirely applicable to these alloys.

[0004] Airbus has developed a technology tailored for laser additive manufacturing. An alloy with the composition Al-4.5Mg-0.51Mn-0.66Sc-0.37Zr was obtained by utilizing the rapid solidification characteristics of laser powder bed melting to obtain a hypereutectic aluminum-scandium modified aluminum alloy with a yield strength exceeding 500 MPa and a tensile strength exceeding 520 MPa. However, the high cost of Sc hinders mass production in engineering. Therefore, researchers have focused on L-PBF forming. Research on alloy modification.

[0005] Therefore, there is a need to develop a high-strength aluminum alloy suitable for the L-PBF process and to customize a special heat treatment process for it in order to achieve excellent mechanical properties to meet the needs of engineering applications. Summary of the Invention

[0006] The technical problem to be solved by this invention is to address the shortcomings of the prior art by providing a high-strength aluminum alloy suitable for laser powder bed melting manufacturing. This aluminum alloy completely replaces the expensive Sc with Hf, and combined with the rapid solidification characteristics of the laser powder bed melting process, the Al3(Zr,Hf) phase formed by Zr, Hf, and the Al matrix acts as a nucleating agent, promoting significant grain refinement in the aluminum alloy and enhancing grain boundary strengthening, thereby increasing the strength of the aluminum alloy and solving the problem of high cost in existing high-strength aluminum alloys.

[0007] To solve the above-mentioned technical problems, the technical solution adopted by the present invention is as follows: a high-strength aluminum alloy suitable for laser powder bed melting manufacturing, characterized in that the aluminum alloy is composed of the following components by mass percentage: Mn 5.50%~6.50%, Mg 2.10%~2.50%, Zr 0.50%~0.90%, Hf 0.90%~1.60%, with the balance being aluminum. According to the designed composition of the target product aluminum alloy, Al-Mn-Mg-Zr-Hf aluminum alloy powder is prepared by vacuum induction gas atomization and then formed by laser powder bed melting. The aluminum alloy has a room temperature tensile strength of ≥520MPa, a yield strength of ≥460MPa, and an elongation after fracture of ≥12% in the deposited state, and a tensile strength of ≥570MPa, a yield strength of ≥520MPa, and an elongation after fracture of ≥5% in the heat-treated state.

[0008] The Al-Mn-Mg-Zr-Hf alloy designed and developed in this invention firstly completely replaces Sc with Hf. Because Hf has a high solid solubility under rapid solidification conditions, exceeding the equilibrium solid solubility limit (0.18%), and exhibits a low equilibrium solid solubility and slow diffusion rate at the commonly used aging temperature of 300℃ for aluminum alloys, Hf forms an Al3Hf phase with the Al matrix. This Al3Hf phase has a smaller lattice mismatch with other phases such as Al3Zr, and possesses better thermal stability and a higher shear modulus. Aluminum alloys can maintain high strength even at high temperatures. Increasing the Zr content allows Zr and Hf to form the Al3(Zr,Hf) phase with the Al matrix. The abundant primary Al3(Zr,Hf) phases significantly refine the grains in the aluminum alloy by increasing nucleation sites, forming high-density grain boundaries that effectively hinder dislocation movement. This effectively improves the strength of the aluminum alloy and enhances its resistance to coarsening. It also reduces lattice mismatch with Al, weakens the coarsening effect of precipitates, and enhances the strengthening effect of the second phase, further improving the strength of the aluminum alloy.

[0009] Compared to The aluminum alloy of this invention incorporates more Mn. Mn acts as a solid solution strengthening agent after rapid solidification and a precipitation strengthening agent after heat treatment, exerting a solid solution strengthening effect during the solidification process and after heat treatment, thereby improving the strength of the aluminum alloy. Simultaneously, it enhances the thermal stability of the Al3X precipitate phase formed by other elements with the Al matrix. Furthermore, the large amount of Mg solute atoms added to the aluminum alloy also dissolves in the Al matrix, exerting a solid solution strengthening effect. The resulting local strain field interacts with dislocations, preventing free dislocation movement and further improving the strength of the aluminum alloy.

[0010] Furthermore, because the elastic modulus of the alloying elements added to the aluminum alloy of this invention is higher than that of the Al matrix, the elastic modulus of the aluminum alloy is increased, resulting in better stiffness. The high-strength aluminum alloy of this invention has an elastic modulus exceeding 75 GPa in the deposited state and an elastic modulus exceeding 80 GPa in the heat-treated state.

[0011] The above-mentioned high-strength aluminum alloy suitable for laser powder bed melting manufacturing is characterized in that the aluminum alloy is composed of the following components by mass percentage: Mn 6.14%, Mg 2.12%, Zr 0.90%, Hf 1.15%, with the balance being aluminum.

[0012] The aforementioned high-strength aluminum alloy suitable for laser powder bed melting manufacturing is characterized in that the particle size of the Al-Mn-Mg-Zr-Hf alloy powder is 20μm to 70μm. This invention utilizes the aforementioned Al-Mn-Mg-Zr-Hf alloy powder with a concentrated particle size distribution and uniform particle size, reducing powder spheroidization and agglomeration during the laser powder bed melting process, and achieving a higher surface finish, thus ensuring the consistency and uniformity of the laser powder bed melting process.

[0013] The above-mentioned high-strength aluminum alloy suitable for laser powder bed melting manufacturing is characterized in that the forming parameters of the laser powder bed melting process are: laser power 350W, laser scanning speed 1200mm / s, scanning spacing 120μm, powder layer thickness 30μm, and interlayer rotation angle 67°.

[0014] The aforementioned high-strength aluminum alloy suitable for laser powder bed melting manufacturing is characterized by an aging heat treatment regime of 375°C for 4 hours followed by air cooling. This invention, by controlling the aging heat treatment regime of the aluminum alloy, ensures the formation of a large number of nano / submicron rod-shaped Al6Mn precipitates and Al3(Zr,Hf) nano-precipitates with an average diameter of 2.5nm ± 0.4nm, thereby improving the peak hardness of the aluminum alloy in the deposited state. Furthermore, this aging heat treatment regime only requires a moderate temperature, making it simpler and easier to implement.

[0015] The aforementioned high-strength aluminum alloy suitable for laser powder bed melting manufacturing is characterized by having a multi-level heterogeneous microstructure: the microstructure at the center of the molten pool consists of micron-sized coarse columnar grains (CCG) with an average grain size of 2.83 μm ± 1.57 μm, and medium-sized equal-axis grains (MEG) with an average grain size of 1.94 μm ± 1.36 μm; the microstructure at the molten pool boundary is mainly submicron-sized ultrafine equiaxed grains (UFG) with an average grain size of 0.56 μm ± 0.01 μm. During the laser powder bed melting process of this aluminum alloy, the higher cooling rate at the molten pool boundary easily forms ultrafine equiaxed grains (UFG), while the lower cooling rate inside the molten pool easily forms coarse columnar grains (CCG), along with medium-sized equal-axis grains (MEG) in between. The process involves the addition of Zr and Hf to the aluminum alloy to form submicron-sized primary Al3(Zr,Hf) phases with the Al matrix, which are distributed in the center of the ultrafine equiaxed grain region. The lattice parameters of the L12 structure Al3(Zr,Hf) are 0.4126 nm ± 0.0104 nm, with low mismatch and good coherence with the α-Al matrix, thus providing non-uniform nucleation sites for the α-Al matrix, promoting grain refinement and enhancing grain boundary strengthening. At the same time, the Mn phase is uniformly distributed in the aluminum alloy as irregular blocky and elongated strip particles, precipitating along the grain boundaries, playing a precipitation strengthening role, thereby improving the strength of the aluminum alloy.

[0016] The preparation process of the high-strength aluminum alloy of this invention is as follows:

[0017] (1) According to the design composition of the target product aluminum alloy, aluminum alloy powder was prepared by vacuum induction gas atomization method, and the composition of aluminum alloy powder was detected by inductively coupled plasma optical emission spectrometer.

[0018] (2) Select aluminum alloy powder of a certain particle size and dry it in a vacuum drying oven at 120°C for 3 hours. Then, use Magic software to model and slice the aluminum alloy. Then, use the German SLM-280HL equipment to perform laser powder bed melting process. The equipment adopts dual laser configuration. The laser wavelength of the fiber laser is 1070nm and the laser spot diameter is 80μm. Before forming, the Al6061 substrate is sandblasted to reduce the surface roughness. Then, it is heated to 150°C and powder is spread and melted to form. The forming parameters are laser power 350W, laser scanning speed 1200mm / s, scanning spacing 120μm, powder layer thickness 30μm, interlayer rotation angle 67°. High-purity argon is used as protective gas during the forming process. The forming process is completed in a glove box with an oxygen content of less than 600ppm.

[0019] Compared with the prior art, the present invention has the following advantages:

[0020] 1. The Al-Mn-Mg-Zr-Hf aluminum alloy of this invention completely replaces the expensive Sc with Hf. Combined with the rapid solidification performance of laser powder bed melting process, the Al3(Zr,Hf) phase formed by Zr, Hf and Al matrix acts as a nucleating agent. By increasing the nucleation sites and promoting heterogeneous nucleation, it promotes significant grain refinement in aluminum alloy, improves grain boundary strengthening and precipitation strengthening effects, thereby improving the strength of aluminum alloy. At the same time, it significantly reduces the cost of aluminum alloy raw materials, solves the problem of high cost of existing aluminum alloys, and provides a high-strength and cost-controllable special aluminum alloy for laser powder bed melting manufacturing, which is suitable for the aerospace field.

[0021] 2. The aluminum alloy of the present invention effectively enhances the strength of the aluminum alloy by adding a large amount of solute element Mn with high equilibrium solubility and high extended solubility under rapid solidification conditions to the Al matrix, which serves as a solid solution reinforcing agent after rapid solidification and a precipitation reinforcing agent after heat treatment.

[0022] 3. This invention produces nearly fully dense aluminum alloy forming parts through laser powder bed melting (L-PBF) process. The aluminum alloy has excellent processing performance, no solidification cracks or obvious metallurgical defects, and is suitable for preparing parts with complex structures.

[0023] 4. This invention employs a simpler medium-temperature aging heat treatment process compared to other additive manufacturing aluminum alloys. This process allows alloying elements with low equilibrium solid solubility in Al to produce fine, dispersed precipitates after heat treatment, thereby enhancing the strength of the aluminum alloy.

[0024] 5. The aluminum alloy prepared by the present invention through laser powder bed melting technology has a multi-level non-uniform heterogeneous structure of micron-level coarse columnar crystals (CCG), medium-equiaxed crystals (MEG), and submicron-level ultrafine equiaxed crystals (UFG). The primary Al3(Zr,Hf) phase distributed in the center of the ultrafine equiaxed crystal region provides non-uniform nucleation sites for the Al matrix, promotes grain refinement, enhances grain boundary strengthening, and thus improves the strength of the aluminum alloy.

[0025] The technical solution of the present invention will be further described in detail below with reference to the accompanying drawings and embodiments. Attached Figure Description

[0026] Figure 1a The image shows the morphology of the aluminum alloy powder prepared in Example 1 of this invention (1000×).

[0027] Figure 1b The image shows the morphology of the aluminum alloy powder prepared in Example 1 of this invention (10000×).

[0028] Figure 2 This is a microstructure diagram of the aluminum alloy in the deposited state prepared in Example 1 of the present invention.

[0029] Figure 3a This is an EBSD image of the aluminum alloy in the deposited state prepared in Example 1 of the present invention.

[0030] Figure 3b This is an EBSD image of the aluminum alloy in the heat-treated state prepared in Example 1 of the present invention.

[0031] Figure 4a The image shows the engineering stress-strain curve of the aluminum alloy prepared in Example 1 of this invention.

[0032] Figure 4b This is a bar chart showing the tensile properties of the aluminum alloy prepared in Example 1 of the present invention.

[0033] Figure 5 This is a STEM-EDS image of the aluminum alloy in the deposited state prepared in Example 1 of the present invention.

[0034] Figure 6a The HAADF diagram and corresponding EDS diagram of the ultrafine grain region in the deposited state of the aluminum alloy prepared in Example 1 of this invention are shown.

[0035] Figure 6b for Figure 6a A magnified view of the area within the dashed line.

[0036] Figure 6c The images shown are HRTEM images and corresponding FFT images of the aluminum alloy in the deposited state prepared in Example 1 of this invention. Detailed Implementation

[0037] Example 1

[0038] The high-strength aluminum alloy of this embodiment is composed of the following components by mass percentage: Mn 6.14%, Mg 2.12%, Zr 0.90%, Hf 1.15%, with the balance being aluminum.

[0039] The method for preparing the high-strength aluminum alloy in this embodiment includes the following steps:

[0040] Step 1: According to the designed composition of the target product aluminum alloy, aluminum alloy powder was prepared using vacuum induction gas atomization. The composition of the aluminum alloy powder was analyzed using inductively coupled plasma optical emission spectrometry (ICP-OES). The results were: Mn 6.10%, Mg 2.22%, Zr 0.81%, Hf 1.09%, with the balance being aluminum. The particle size of the aluminum alloy powder was analyzed according to GB / T-19077 "Laser Diffraction Method for Particle Size Distribution". The results were: D10 = 20.1 μm, D50 = 38.2 μm, D90 = 64.8 μm, and the powder exhibited good sphericity. Figure 1a and Figure 1b As shown, most of the aluminum alloy powder is spherical, with only a small amount of satellite powder present;

[0041] (2) Aluminum alloy powder with a particle size of 20μm to 70μm was placed in a vacuum drying oven and dried at 120℃ for 3 hours to remove moisture. Then, the aluminum alloy was modeled and sliced ​​using Magic software. A test block was then formed using a laser powder bed melting process with a German SLM-280HL device. This device uses a dual-laser configuration, with the fiber laser having a wavelength of 1070nm and a laser spot diameter of 80μm. Before forming, the Al6061 substrate was sandblasted to reduce [the risk of damage]. The surface roughness was then determined, followed by heating to 150℃, powder spreading, and melting to form the desired shape. The forming parameters were: laser power 350W, laser scanning speed 1200mm / s, scanning spacing 120μm, powder layer thickness 30μm, and interlayer rotation angle 67°. High-purity argon was used as a protective gas during the forming process, which was completed in a glove box with an oxygen content of less than 600ppm. Finally, wire cutting was used to separate the sample from the substrate to obtain a cubic aluminum alloy with dimensions of 10mm×10mm×10mm, i.e., deposited aluminum alloy.

[0042] Testing showed that the density of the deposited aluminum alloy prepared in this embodiment was above 99.7%. SEM samples were prepared using standard metallographic methods and observed under a scanning electron microscope. The results are as follows: Figure 2 As shown.

[0043] Figure 2 This is a microstructure image of the deposited aluminum alloy prepared in this embodiment. Figure 2 It can be seen that the aluminum alloy has a multi-level non-uniform heterogeneous structure in the deposited state: the structure in the center of the molten pool is micron-sized coarse columnar crystal CCG (at b1) and medium-sized isoaxial crystal MEG (at b2), while the structure at the molten pool boundary (BD) is mainly submicron-sized ultrafine isoaxial crystal UFG (at b3).

[0044] The deposited aluminum alloy prepared in this embodiment was mechanically polished and then subjected to stress relief treatment using a Leica ion polisher. EBSD testing was performed under a focused ion / electron dual-beam system, and the results are as follows: Figure 3a As shown.

[0045] Figure 3a The image shown is an EBSD image of the aluminum alloy in the deposited state prepared in this embodiment. Figure 3a It can be seen that the micron-sized coarse columnar crystals (CCG) and medium-sized isometric crystals (MEG) are located inside the molten pool, while the submicron-sized ultrafine isometric crystals (UFG) are located at the bottom of the molten pool. Based on the EBSD plot, the area fraction of the CCG and MEG regions inside the molten pool is approximately 43%, while the area fraction of the UFG region at the bottom of the molten pool is approximately 57%.

[0046] The deposited aluminum alloy prepared in this embodiment was held at 375℃ for 4 hours and then air-cooled for aging heat treatment to obtain a heat-treated aluminum alloy. Simultaneously, the deposited aluminum alloy prepared in this embodiment was held at 375℃ for 3 hours, 8 hours, and 10 hours respectively as a control heat-treated aluminum alloy. According to GB / T 228-2002 "Metallic Materials - Tensile Testing at Room Temperature", specimens were prepared from the deposited and heat-treated aluminum alloys and the control heat-treated aluminum alloy, respectively. Then, room temperature tensile tests were performed on an INSTRON 3382 electronic universal testing machine at a displacement rate of 1 mm / min. The results are as follows. Figure 4a and Figure 4b As shown, the results indicate that the stress-strain curves of aluminum alloys heat-treated at 375℃ for different times have high overlap, and their yield strength and tensile strength are close, while their elongation has slight differences. Among them, the heat-treated aluminum alloy obtained by heat treatment at 375℃ for 4 hours has the best comprehensive mechanical properties. Compared with the aluminum alloy in the deposited state, which has a room temperature tensile strength of 520MPa, a yield strength of 464MPa, an elongation after fracture of 12%, and an elastic modulus of 78GPa, the heat-treated aluminum alloy has a tensile strength of 578MPa, a yield strength of 525MPa, an elongation after fracture of 5%, and an elastic modulus of 83GPa.

[0047] Figure 3b The EBSD image of the aluminum alloy in the heat-treated state prepared in this embodiment shows that the area fraction of the CCG region and UFG region inside the molten pool is about 43%, and the area fraction of the UFG region at the bottom of the molten pool is about 57%.

[0048] Will Figure 3a and Figure 3b The comparison shows that the grain size of the aluminum alloy in the heat-treated state did not change significantly compared with the deposited state, indicating that the grains of the aluminum alloy did not grow under the heat treatment conditions.

[0049] Figure 5 The image shown is a STEM-EDS image of the aluminum alloy in the deposited state prepared in this embodiment. Figure 5 It can be seen that the Mn and Mg phases are uniformly distributed in the aluminum alloy. The Mn phase precipitates as granular and strip-shaped precipitates along the grain boundaries, and the Mg phase also segregates at the grain boundaries. In contrast, Zr and Hf elements only precipitate significantly in the ultrafine grain region, and the precipitate size is relatively small.

[0050] Figure 6a The images show the HAADF diagram and corresponding EDS diagram of the ultrafine-grained region in the deposited state of the aluminum alloy prepared in this embodiment. Figure 6b for Figure 6a A magnified view of the area within the dashed line. Figure 6cThe HRTEM image and corresponding FFT image of the aluminum alloy in the deposited state prepared in this embodiment are shown in the figure. Figures 6a-6c It can be seen that cubic precipitates enriched with Zr and Hf were observed in the grain center of the ultrafine crystal region, with an average size of about 87 nm ± 29 nm. The diffraction spots can be identified as Al3(Zr,Hf) phase.

[0051] Example 2

[0052] The high-strength aluminum alloy of this embodiment is composed of the following components by mass percentage: Mn 6.46%, Mg 2.35%, Zr 0.60%, Hf 1.58%, with the balance being aluminum.

[0053] The difference between the preparation method of the high-strength aluminum alloy in this embodiment and that in Embodiment 1 is that the laser powder bed melting process parameters are: laser power 325W, laser scanning speed 1000mm / s, and finally wire cutting is used to separate the test block from the substrate to obtain a cubic aluminum alloy with dimensions of 10mm×10mm×10mm, i.e., the deposited aluminum alloy; then the deposited aluminum alloy is held at 375℃ for 4h and air-cooled for aging heat treatment to obtain the heat-treated aluminum alloy.

[0054] According to GB / T 228-2002 "Metallic Materials - Tensile Testing at Room Temperature", deposited and heat-treated aluminum alloys prepared in this embodiment were used to prepare deposited and heat-treated specimens, respectively. Room temperature tensile tests were then conducted on an INSTRON 3382 electronic universal testing machine at a displacement rate of 1 mm / min. The results showed that the room temperature tensile strength of the deposited aluminum alloy was 520 MPa, the yield strength was 472 MPa, the elongation after fracture was 12%, and the elastic modulus was 79 GPa. The tensile strength of the heat-treated aluminum alloy was 584 MPa, the yield strength was 525 MPa, the elongation after fracture was 5%, and the elastic modulus was 82 GPa.

[0055] Example 3

[0056] The high-strength aluminum alloy of this embodiment is composed of the following components by mass percentage: Mn 5.56%, Mg 2.48%, Zr 0.51%, Hf 0.94%, with the balance being aluminum.

[0057] The difference between the preparation method of the high-strength aluminum alloy in this embodiment and that in Embodiment 1 is that the laser powder bed melting process parameters are: laser power 325W, laser scanning speed 1200mm / s, and finally wire cutting is used to separate the test block from the substrate to obtain a cubic aluminum alloy with a size of 10mm×10mm×10mm, i.e., the deposited aluminum alloy; then the deposited aluminum alloy is held at 375℃ for 4h and air-cooled for aging heat treatment to obtain the heat-treated aluminum alloy.

[0058] According to GB / T 228-2002 "Metallic Materials - Tensile Testing at Room Temperature", deposited and heat-treated aluminum alloys prepared in this embodiment were used to prepare deposited and heat-treated specimens, respectively. Room temperature tensile tests were then conducted on an INSTRON 3382 electronic universal testing machine at a displacement rate of 1 mm / min. The results showed that the room temperature tensile strength of the deposited aluminum alloy was 524 MPa, the yield strength was 475 MPa, the elongation after fracture was 12%, and the elastic modulus was 78 GPa. The tensile strength of the heat-treated aluminum alloy was 581 MPa, the yield strength was 520 MPa, the elongation after fracture was 5%, and the elastic modulus was 81 GPa.

[0059] The above description is merely a preferred embodiment of the present invention and is not intended to limit the invention in any way. Any simple modifications, alterations, and equivalent changes made to the above embodiments based on the inventive essence shall still fall within the protection scope of the present invention.

Claims

1. A high-strength aluminum alloy suitable for laser powder bed melting manufacturing, characterized in that, This aluminum alloy is composed of the following components by mass percentage: Mn 5.50%~6.50%, Mg 2.10%~2.50%, Zr 0.50%~0.90%, Hf The aluminum alloy comprises 0.90%~1.60% aluminum, with the balance being aluminum. Based on the designed composition of the target product aluminum alloy, Al-Mn-Mg-Zr-Hf aluminum alloy powder is prepared using vacuum induction gas atomization and then formed by laser powder bed melting. The aluminum alloy exhibits a room temperature tensile strength of over 520 MPa, a yield strength of over 460 MPa, and an elongation after fracture of over 12% in the deposited state. In the heat-treated state, the tensile strength is over 570 MPa, the yield strength is over 520 MPa, and the elongation after fracture is over 5%. The aluminum alloy possesses a multi-level heterogeneous microstructure: the central microstructure of the molten pool consists of micron-sized coarse columnar crystals (CCG) with an average grain size of 2.83 μm ± 1.57 μm, and medium-sized isoaxial crystals (MEG) with an average grain size of 1.94 μm ± 1.36 μm. The boundary microstructure of the molten pool mainly consists of submicron-sized ultrafine isoaxial crystals (UFG) with an average grain size of 0.56 μm ± 0.01 μm.

2. The high-strength aluminum alloy suitable for laser powder bed melting manufacturing according to claim 1, characterized in that, This aluminum alloy consists of the following components by mass percentage. Composition: Mn 6.14%, Mg 2.12%, Zr 0.90%, Hf 1.15%, balance aluminum.

3. A high-strength aluminum alloy suitable for laser powder bed melting manufacturing according to claim 1, characterized in that, The particle size of the Al-Mn-Mg-Zr-Hf aluminum alloy powder is 20μm~70μm.

4. A high-strength aluminum alloy suitable for laser powder bed melting manufacturing according to claim 1, characterized in that, The forming parameters of the laser powder bed melting process are: laser power 350W, laser scanning speed 1200mm / s, scanning spacing 120μm, powder layer thickness 30μm, and interlayer rotation angle 67°.

5. A high-strength aluminum alloy suitable for laser powder bed melting manufacturing according to claim 1, characterized in that, The aging heat treatment regime adopted for the aluminum alloy in the heat-treated state is: holding at 375℃ for 4 hours, followed by air cooling.

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