High-strength and high-plasticity 7xxx aluminum alloy and preparation method thereof

Abstract

The invention relates to the technical field of additive manufacturing, in particular to a 7xxx aluminum alloy with high strength and plasticity and a preparation method thereof.The preparation method comprises the steps that S1, the surfaces of a 7A52 aluminum alloy bar and an aluminum alloy base plate are cleaned; s2, the cleaned 7A52 aluminum alloy bar serves as a deposition raw material, layer-by-layer deposition is conducted on the cleaned aluminum alloy base plate in a stirring friction deposition additive mode, and a 7A52 aluminum alloy additive body is obtained; and S3, the 7A52 aluminum alloy additive body is sequentially subjected to solution treatment and two-stage aging treatment, and the finished high-strength and high-plasticity 7A52 aluminum alloy is obtained. According to the method, on the premise that alloy components are not changed, through solute homogenization in the solid solution stage and staged precipitation control in the two-stage aging stage, synergistic improvement of the strength and plasticity of the additive body is achieved, and the performance difference value of different deposition layers is remarkably reduced.

Description

Technical Field

[0001] This invention relates to the field of additive manufacturing technology, specifically to a high-strength and ductile 7xxx aluminum alloy and its preparation method. Background Technology

[0002] 7A52 aluminum alloy, as a high-strength aluminum alloy in the Al-Zn-Mg-Cu system, has the characteristics of low density and high specific strength, and has good application prospects in lightweight load-bearing components. Currently, most 7A52 aluminum alloy components are manufactured using a combination of casting and thermomechanical processing, which typically suffers from long manufacturing processes, low material utilization, and high energy consumption. Furthermore, limitations imposed by molds and deformation paths make it difficult to achieve near-net-shape forming and rapid iterative manufacturing of complex components, thus failing to meet the requirements of high-end equipment for rapid manufacturing and green, low-carbon production.

[0003] To improve manufacturing efficiency and reduce material waste, additive manufacturing technology, with its advantages of rapid prototyping and high material utilization, is considered a potentially effective approach for preparing 7A52 aluminum alloy components. However, 7A52 and other 7xxx series aluminum alloys are characterized by high crack sensitivity and poor weldability. In traditional melt additive manufacturing processes (such as laser powder bed melting and arc additive manufacturing), metallurgical defects such as porosity and hot cracks are easily generated due to the solidification shrinkage of the molten pool and the effects of thermal stress. This makes it difficult to achieve both strength and plasticity in additive components, and the mechanical properties often fail to reach the level of similar thermomechanically processed materials. To avoid liquid-solid phase transformation defects in melt additive manufacturing, existing technologies have disclosed solid-state additive manufacturing methods. Among them, additive friction stir deposition (AFSD) induces viscoplastic flow and layer-by-layer deposition of materials through frictional heating and extrusion shearing, which can reduce the risk of defects such as porosity and hot cracks to a certain extent, providing a new process route for the additive manufacturing of 7xxx series aluminum alloys.

[0004] Despite the advantages of solid-state deposition, AFSD (Alternating Aspect Deposition) has shown that its forming quality and microstructure are highly sensitive to heat input, resulting in a relatively narrow forming process window. Particularly in multi-layer deposition processes, repeated thermal cycles make heat accumulation difficult to avoid. Differences in thermal history along the deposition height / thickness direction can lead to variations in the morphology and size of precipitates in different layers, potentially accompanied by microstructural features such as absence of precipitates. This can cause interlayer gradients in performance indicators such as hardness, ultimate tensile strength, and elongation, resulting in insufficient microstructure and property uniformity. To address these issues, existing technologies typically achieve heat input matching by adjusting process parameters such as spindle speed, travel speed, axial load, and overlap method, aiming to simultaneously meet requirements for plastic flow, dynamic recrystallization, and avoidance of localized overheating. Other technologies mitigate interlayer differences by reducing deposition temperature rise or optimizing the deposition path. However, due to the inherent thermal accumulation effect and the thermodynamic and kinetic characteristics of precipitate coarsening, without changing the alloy composition, it is often difficult to fundamentally solve the problems of precipitate coarsening and interlayer inhomogeneity by simply optimizing the AFSD process parameters, resulting in the overall strength and plasticity of the additive body still being difficult to improve in a synergistic way.

[0005] To improve the characteristics and mechanical properties of precipitated phases, existing technologies have disclosed the traditional T6 heat treatment method of solution treatment followed by artificial aging for 7xxx series aluminum alloys. This method aims to improve the alloy strength by dissolving coarse η phases through solution treatment and forming nanoscale η′ phases during the aging stage. However, for fine-grained structures obtained through friction stir processing or solid-state additive manufacturing, conventional T6 heat treatment under solution heat exposure conditions may significantly increase the driving force for grain boundary migration, making the fine-grained structure prone to abnormal grain growth. At the same time, in the absence of pre-nucleation and effective pinning, subsequent artificial aging is prone to non-uniform precipitation along grain boundaries and defects, accompanied by secondary coarsening of the precipitated phases or over-aging, leading to deterioration of hardness and tensile strength. This results in a mismatch between strength and plasticity, further restricting the engineering application of AFSD in the preparation of 7A52 and similar 7xxx series aluminum alloys.

[0006] In summary, the existing technology has at least the following shortcomings: (1) 7A52 aluminum alloy is prepared by casting and thermomechanical processing. The manufacturing cycle is long and the material utilization rate is low. It is also limited by mold and deformation path, which is not conducive to near-net-shape forming and rapid iterative manufacturing of complex components. (2) Traditional melt additive manufacturing of 7xxx series aluminum alloys is prone to defects such as porosity and hot cracking, making it difficult to balance strength and plasticity; (3) Although AFSD can reduce melting defects, the heat accumulation effect of multilayer deposition is difficult to avoid, which easily leads to coarsening of precipitates and interlayer differences, resulting in insufficient uniformity of microstructure and properties. (4) Although conventional T6 heat treatment can improve strength, it poses a risk of abnormal grain growth and coarsening of precipitates in fine-grained additive structures, which can easily lead to a mismatch between strength and plasticity and make it difficult to achieve a synergistic improvement in strength and plasticity.

[0007] Therefore, there is an urgent need to provide a strategy suitable for preparing high-strength and ductile 7A52 aluminum alloys using AFSD+ heat treatment. Summary of the Invention

[0008] To address the problems in the prior art, this invention provides a high-strength and ductile 7xxx aluminum alloy and its preparation method. Without changing the alloy composition, it can achieve homogenization of additive body structure and synergistic improvement of strength and ductility through solute homogenization in the solid solution stage and staged precipitation control in the two-stage aging stage, and significantly reduce the performance difference between different deposition layers.

[0009] This invention is achieved through the following technical solution: A method for preparing high-strength and high-ductility 7xxx series aluminum alloys, comprising: S1, Clean the surface of the 7A52 aluminum alloy rod and aluminum alloy substrate; S2, using the cleaned 7A52 aluminum alloy rod as the deposition material, the 7A52 aluminum alloy additive body is obtained by layer-by-layer deposition on the cleaned aluminum alloy substrate using additive friction stir deposition. S3, the 7A52 aluminum alloy additive body is subjected to solution treatment and two-stage aging treatment in sequence to obtain the finished high-strength and ductile 7A52 aluminum alloy.

[0010] Preferably, in S1, the chemical composition of the 7A52 aluminum alloy bar by mass percentage is: Zn 4.0%~5.0%, Mg 1.8%~2.5%, Cu 0.1%~0.4%, Mn 0.2%~0.5%, Cr 0.1%~0.3%, Al 91.0%~93%, with the remainder being impurities.

[0011] Preferably, in S2, the parameters for layer-by-layer deposition are: axial load of 20–40 kN, spindle speed of 800–1200 rpm, travel speed of 10–40 mm / min, feed speed of 3–8 mm / min, and shoulder diameter of the stirring head of 55–65 mm.

[0012] Preferably, in S2, the deposited microstructure of the 7A52 aluminum alloy additive is a fine equiaxed crystal structure with an average grain size ≤7 μm.

[0013] Preferably, in S3, during the solution treatment, the 7A52 aluminum alloy additive body is heated to 450-490°C and held for 1-3 hours, and then quenched in water at 10-40°C to obtain an intermediate aluminum alloy.

[0014] Preferably, the intermediate aluminum alloy retains Fe and Mn enriched dispersed phases, including Al6(Fe, Mn) phase, with a content of 2.31%~3.36%.

[0015] Preferably, in S3, the two-stage aging treatment involves sequentially performing a first-stage aging treatment and a second-stage aging treatment on the solution-treated 7A52 aluminum alloy additive body. During the first-stage aging treatment, the temperature is 90–115 ℃ and the holding time is 4–10 h to obtain crude aluminum alloy. During the secondary aging treatment, the temperature is 120-150 ℃ and the holding time is 10-20 h to obtain a finished high-strength and ductile 7A52 aluminum alloy with η′ precipitates.

[0016] Preferably, in S3, the equivalent size of the η′ precipitate is 15–30 nm.

[0017] Preferably, the finished high-strength and high-ductility 7A52 aluminum alloy is made of Al. 18 The Mg3Cr2 phase contains a nanotwin structure.

[0018] A high-strength and ductile aluminum alloy obtained by a high-strength and ductile 7xxx series aluminum alloy preparation method, wherein the difference in ultimate tensile strength of different deposited layers along the deposition height direction is ≤30 MPa and the difference in elongation is ≤6%.

[0019] Compared with the prior art, the present invention has the following beneficial effects: This invention discloses a method for preparing high-strength and ductile 7xxx series aluminum alloys. The method employs AFSD solid-state deposition to prepare 7A52 aluminum alloy additive bodies, reducing the risk of defects such as porosity and hot cracking in traditional melt additive manufacturing. It also offers advantages such as high material utilization, strong near-net-shape forming capability, and suitability for rapid iterative manufacturing of complex components. Without altering the alloy composition, a staged control strategy involving solution treatment and two-stage aging treatment is used to regulate the nucleation and growth behavior of precipitates. This improves the refinement and spatial uniformity of the precipitates while suppressing abnormal growth of fine-grained structures and the risk of precipitate re-coarsening during solution / aging. This enhances the consistency of microstructure and properties across different deposition layers, achieving a synergistic improvement in the strength and ductility of the 7A52 aluminum alloy.

[0020] Furthermore, strictly limiting parameters such as deposition rate, thickness, and temperature for each layer ensures that the product size, shape, and microstructure closely match the target product, reducing the accumulation of interlayer errors. The shoulder diameter of the stirring head is 55–65 mm to ensure the formation of the 7A52 aluminum alloy additive body. This shoulder diameter range provides a suitable frictional contact area and axial forging action, promoting full plasticization, continuous spreading, and good interlayer bonding of the deposited material. If the shoulder diameter is too small, insufficient heat input and forging coverage will easily lead to insufficient material plasticization, limited deposition channel width, and discontinuous forming; if the shoulder diameter is too large, excessive heat input, increased heat accumulation, and increased flash will easily occur, which is detrimental to dimensional control and maintaining microstructure uniformity. Therefore, controlling the shoulder diameter within the range of 55–65 mm is beneficial for balancing the forming stability, controllable heat input, and consistent microstructure properties during the 7A52 aluminum alloy AFSD process.

[0021] Furthermore, limiting the grain size of the fine equiaxed crystal structure in the deposited microstructure of the 7A52 aluminum alloy additive is to ensure the strengthening effect of fine grains, which helps to improve the material strength and plasticity, thus achieving synergistic optimization of strength and plasticity. If the grain size is too large, the strengthening effect will be weakened.

[0022] Furthermore, solution treatment allows the coarse precipitates to be re-dissolved and solutes to be redistributed, providing a homogeneous solute basis for subsequent precipitation.

[0023] Furthermore, the Fe and Mn enriched dispersed phases retained after solution treatment can provide pinning effect on grain boundary migration, which helps improve the stability of the fine-grained structure during subsequent heat exposure and reduces the risk of abnormal grain growth that easily occurs in conventional single-stage aging. Since fine grains can provide significant grain boundary strengthening, the contribution of grain boundary strengthening to the yield strength in the additive body obtained in this invention can reach 71.7 MPa, accounting for approximately 21% of the total strength. This indicates that the maintenance of grain size stability by Fe and Mn enriched dispersed phases is not only reflected in the retention of fine grains at the microstructure level but can also be directly converted into an improvement in macroscopic mechanical properties. If abnormal grain coarsening occurs during subsequent heat treatment, the number of grain boundaries will decrease, and the contribution of grain boundary strengthening will be significantly weakened, thus hindering the synergistic maintenance of yield strength and ductility. On the other hand, if the dispersed phase content is too high, it will induce crack initiation during tensile testing, impairing tensile properties; if the content is too low, the pinning effect will be lost.

[0024] Furthermore, a two-stage aging process is employed to achieve precise control of the microstructure. Specifically, the first-stage aging, under conditions of high supersaturation and limited diffusion, is used to form high-density, widely distributed GP regions within the grain as nucleation sites, expanding the nucleation sites from local dominant sites such as grain boundaries / dislocations to universal sites within the grain. During the second-stage aging, the solute is preferentially consumed near these pre-set nuclei and grows into the η′ phase under controlled conditions. The diffusion fields of multiple precipitates overlap to suppress local coarsening and uneven growth, thereby reducing the number density and size differences of precipitates in different regions (obtaining fine precipitates with a scale of 15–30 nm) and achieving spatial homogenization.

[0025] Furthermore, the nanotwin structure is formed by the addition material undergoing two-stage aging treatment during tensile deformation in Al... 18 The nanotwin structure formed within the Mg3Cr2 phase is beneficial for achieving better plastic coordination and work hardening during the deformation process, thereby obtaining a 7A52 aluminum alloy additive body with both high strength and high plasticity. In preferred cases, the ultimate tensile strength and elongation of the finished high-strength and high-plasticity 7A52 aluminum alloy reach 99% and 275% of those of conventional thermomechanical alloys, respectively.

[0026] This invention discloses a method for preparing 7A52 aluminum alloy using a high-strength and ductile 7xxx series aluminum alloy. Through staged precipitation control during two-stage aging, the difference in precipitation between layers and the performance gradient caused by the thermal accumulation effect of AFSD multilayer deposition can be reduced. This results in a difference in ultimate tensile strength between different deposition layers of ≤30 MPa and a difference in elongation of ≤6%, thereby improving the overall reliability of the component. In other words, the process proposed in this invention can maintain the strength of aluminum alloys obtained by traditional preparation methods while greatly improving elongation, i.e., plastic deformation capacity. Attached Figure Description

[0027] Figure 1 This is a schematic flowchart of a method for preparing high-strength and ductile 7xxx series aluminum alloys according to the present invention.

[0028] Figure 2 This is a schematic diagram of the AFSD deposition path in the preparation method of a high-strength and ductile 7xxx series aluminum alloy of the present invention.

[0029] Figure 3 A schematic diagram of the sampling locations for additive manufacturing (sampling along the deposition height direction at the top / middle / bottom layers, and a schematic diagram of the orientation of the tensile specimens). Figure 4 The images shown are of the deposited 7A52 aluminum alloy additive body according to Embodiment 1 of the present invention. (a) is a surface view of the 7A52 aluminum alloy additive body, (b) is a cross-sectional view of the deposited 7A52 aluminum alloy additive body, and (c) is a side view of the deposited 7A52 aluminum alloy additive body.

[0030] Figure 5 The images show the EBSD orientation diagram and grain boundary distribution diagram of the deposited 7A52 aluminum alloy additive body in Embodiment 1 of the present invention, where (a, e, i) is the grain orientation diagram, (b, f, j) is the grain size statistics diagram, (c, g, k) is the grain boundary diagram, and (d, h, l) is the phase diagram.

[0031] Figure 6 This is a distribution diagram of the precipitated phases in the 7A52 aluminum alloy additive body after solution treatment in Example 1 of the present invention.

[0032] Figure 7 The image shows the precipitate distribution of 7A52 aluminum alloy after aging treatment in Example 1 of the present invention. In this image, (a) is the precipitate distribution diagram, (bg) is the corresponding EDS energy spectrum, (h) is the phase type analyzed by XRD pattern, and (ik) is the morphology of precipitates in different layers.

[0033] Figure 8 The images shown are TEM images of dispersed phases and precipitated phases near grain boundaries and subgrain boundaries after aging treatment in Embodiment 1 of the present invention. Among them, (a) is a low-magnification grain boundary morphology, (b) is a subgrain boundary pinned by the precipitated phase inside the grain, (c) is an elemental distribution map of the precipitated phase in this region, (d) is a map of MgZn2 phase pinning grain boundaries, (e) is an enlarged view of the junction between the MgZn2 phase and the grain boundary, (f) is a strain map at the junction of the phase and the grain boundary, and (g) is a map of lattice distortion caused by the phase.

[0034] Figure 9 The image shows the intraphase twin diagram of Al18Mg3Cr2 in Example 1 of this invention, where (a) is a morphology diagram of the intraphase twins, (b) is a selected area electron diffraction pattern for determining the type of precipitated phase and the twin relationship, and (cf) is an EDS energy spectrum of the phase.

[0035] Figure 10 This is a tensile curve of the 7A52 aluminum alloy in Embodiment 1 of the present invention.

[0036] Figure 11 This is a stretching curve diagram of Comparative Example 1 of the present invention.

[0037] Figure 12 This is a stretching curve diagram of Comparative Example 2 of the present invention. Detailed Implementation

[0038] The following specific embodiments illustrate the implementation of the present invention. Those skilled in the art can easily understand other advantages and effects of the present invention from the content disclosed in this specification.

[0039] Exemplary embodiments of the present invention will now be described with reference to the accompanying drawings. However, the present invention may be embodied in many different forms and is not limited to the embodiments described herein. These embodiments are provided to fully and completely disclose the invention and to fully convey its scope to those skilled in the art. The terminology used in the exemplary embodiments illustrated in the drawings is not intended to limit the invention. In the drawings, the same units / elements are referred to by the same reference numerals.

[0040] Unless otherwise stated, the terms used herein (including technical terms) have their common meaning as understood by one of ordinary skill in the art. Furthermore, it is understood that terms defined in commonly used dictionaries should be understood to have a meaning consistent with the context of their relevant field, and not to be interpreted as having an idealized or overly formal meaning.

[0041] The present invention will be further described in detail below with reference to specific embodiments. These descriptions are for explanation purposes only and are not intended to limit the scope of the invention.

[0042] This invention discloses a method for preparing high-strength and high-ductility 7xxx series aluminum alloys, referring to... Figure 1 ,include: S1, Clean the surface of the 7A52 aluminum alloy rod and aluminum alloy substrate.

[0043] Among them, 7A52 aluminum alloy rods are selected as the deposition raw materials. The rod diameter is 20~30 mm and the length is 200~500 mm. The chemical composition by mass percentage is: Zn 4.0%~5.0%, Mg 1.8%~2.5%, Cu 0.1%~0.4%, Mn 0.2%~0.5%, Cr 0.1%~0.3%, and the remainder is Al and impurities.

[0044] The aluminum alloy substrate is made of 6061 aluminum alloy with a thickness of 5~20 mm, and is fixed to the worktable by a clamp.

[0045] The end face of the rod and the surface of the substrate are mechanically ground or sandblasted to remove the oxide layer, and then cleaned with anhydrous ethanol / acetone. The substrate is then fixed on the worktable.

[0046] S2, using the cleaned 7A52 aluminum alloy rod as the deposition material, the additive friction stir deposition equipment (equipped with a pressurization system that can realize axial load control, and the diameter of the stirring head shoulder is 55-65 mm) is used to deposit layer by layer on the cleaned aluminum alloy substrate to obtain the 7A52 aluminum alloy additive body.

[0047] The parameters for layer-by-layer deposition are as follows: axial load of 20–40 kN, spindle speed of 800–1200 rpm, travel speed of 10–40 mm / min, and feed speed of 3–8 mm / min.

[0048] The deposited microstructure of the 7A52 aluminum alloy additive is a fine equiaxed crystal structure with an average grain size of ≤7 μm.

[0049] Deposition path reference Figure 2 As shown, after the shoulder reaches the predetermined position, the spindle begins to rotate and feeds material synchronously. Under the action of friction and extrusion, the raw material undergoes viscoplastic flow and fills the gap between the shoulder and the substrate. As the shoulder moves along the predetermined trajectory, the plasticized raw material first forms a metallurgical adhesion with the substrate, and then spreads layer by layer to build the additive body. When the shoulder moves to the end of the trajectory, the feeding stops, the shoulder rises and finally stops rotating, and the deposition process ends.

[0050] S3, using a box-type resistance furnace to sequentially perform solution treatment and two-stage aging treatment on the 7A52 aluminum alloy additive body, to obtain the finished high-strength and ductile 7A52 aluminum alloy.

[0051] During solution treatment, the 7A52 aluminum alloy additive body is heated to 450-490℃ and held for 1-3 hours, and then quenched in water at 10-40℃ to obtain an intermediate aluminum alloy. The intermediate aluminum alloy retains Fe and Mn enriched dispersed phases, including Al6(Fe, Mn) phase with a content of 2.31%-3.36%.

[0052] The two-stage aging treatment involves sequentially performing a first-stage aging treatment and a second-stage aging treatment on the solution-treated 7A52 aluminum alloy additive body. During the first-stage aging treatment, the temperature is 90–115 ℃ and the holding time is 4–10 h to obtain crude aluminum alloy. During the secondary aging treatment, the temperature is 120–150 °C and the holding time is 10–20 h to obtain a finished high-strength and ductile 7A52 aluminum alloy with η′ precipitates, and the equivalent size of the η′ precipitates is 15–30 nm.

[0053] Finished high-strength and ductile 7A52 aluminum alloy Al 18 The Mg3Cr2 phase contains a nanotwin structure.

[0054] The present invention also discloses a high-strength and ductile aluminum alloy obtained by a high-strength and ductile 7xxx series aluminum alloy preparation method, wherein the difference in ultimate tensile strength of different deposited layers along the deposition height direction is ≤30 MPa and the difference in elongation is ≤6%.

[0055] Example 1 S1 is made of 7A52 aluminum alloy rod with a diameter of 20 mm and 6061 aluminum alloy as the base material. S2: Layer-by-layer deposition was performed using AFSD, with the following deposition parameters: axial load of 30 kN, spindle speed of 1000 rpm, travel speed of 20 mm / min, feed rate of 5 mm / min, and agitator shoulder diameter of 60 mm. A parallel three-layer overlapping path was used for layer-by-layer deposition to obtain a 7A52 aluminum alloy additive body. S3, solution treatment of 7A52 aluminum alloy additive body: heat to 470℃ and hold for 2 h, then water quench; The solution-treated additive body was subjected to a two-stage aging treatment: the first-stage aging temperature was 105℃ and the holding time was 8 h; followed by the second-stage aging temperature was 130℃ and the holding time was 14 h, to obtain the finished high-strength and ductile 7A52 aluminum alloy.

[0056] Figure 4 The image shown is of the actual 7A52 aluminum alloy additive body in the deposited state in Example 1. It can be seen that the additive body has no obvious defects and is well formed.

[0057] Figure 5 The images show the EBSD orientation diagram and grain boundary distribution diagram of the deposited 7A52 additive. The top, middle and bottom layers of the additive have small equiaxed grains with random orientations. The average grain sizes of the three layers are 6.03 μm, 6.23 μm and 6.76 μm, respectively.

[0058] Figure 6 The image shows the phase precipitate of the additive body after solution treatment, where the Al6(Fe, Mn) phase can be observed.

[0059] Figure 7 The image shows the precipitate distribution of the additive after aging treatment. Both the GP region and the η′ phase were observed, with an average size of 20 nm.

[0060] Figure 8 The images show TEM images of dispersed and precipitated phases near grain boundaries and subgrain boundaries after aging treatment. Al6(Fe,Mn) phase and MgZn2 phase pinning grain boundaries can be observed.

[0061] Figure 9 Al after stretching deformation 18 The diagrams of twinning, stacking faults, and 9R transition structures within the Mg3Cr2 phase also demonstrate that this embodiment activated a new deformation mechanism, resulting in Al after tensile deformation. 18The twin / stacked fault interface within the Mg3Cr2 phase can absorb local strain energy and alleviate stress concentration at the particle / matrix interface, thereby inhibiting particle cracking and interface debonding, and delaying micropore initiation. On the other hand, it can act as a high-density internal interface, significantly hindering dislocation movement and promoting dislocation storage, thus improving work hardening capacity and delaying necking. The presence of the 9R transition structure further demonstrates the continuity of shear evolution within the phase, which can continuously contribute energy dissipation and hardening over a wider strain range. Therefore, the twin-stacked fault-9R structure evolution within this phase constitutes an auxiliary deformation mechanism distinct from matrix dislocation slip, and its synergistic effect with fine-grained microstructure and η′ nanoprecipitation strengthening contributes to achieving a better strength-plasticity match.

[0062] Figure 10 The tensile curve of the additive body of this invention shows that each layer of the additive body exhibits excellent strength-plasticity matching, with the top layer having UTS, YS, and EL values ​​of 473 MPa, 357 MPa, and 22%, respectively. The ultimate tensile strength and elongation of this 7A52 aluminum alloy reach 99% and 275% of those of conventionally thermomechanically processed 7A52 aluminum alloy (i.e., conventionally rolled 7A52 alloy treated with T6 has a tensile strength of 480 MPa and an elongation of 8%, Y. Feng, J. Chen, W. Qiang, Y. Wang, Z. Li, Y. Zhang, Microstructure and mechanical properties of aluminum alloy 7A52 thick plates welded by robotic double-sided coaxial GTAW process, Mater. Sci. Eng., A 673(2016) 8-15.).

[0063] Example 2: S1 is made of 7A52 aluminum alloy rod with a diameter of 20 mm and 6061 aluminum alloy as the base material.

[0064] S2, using AFSD for layer-by-layer deposition, with deposition parameters meeting the following requirements: axial load of 20 kN, spindle speed of 1200 rpm, travel speed of 40 mm / min, feed rate of 8 mm / min, and agitator shoulder diameter of 55 mm. A parallel three-layer overlapping path was used for layer-by-layer deposition to obtain a 7A52 aluminum alloy additive body. S3, solution treatment of 7A52 aluminum alloy additive body: heat to 450℃ and hold for 3 h, then water quench; The solution-treated additive body was subjected to a two-stage aging treatment: the first stage aging temperature was 90℃ and the holding time was 10 h; followed by the second stage aging temperature at 150℃ and the holding time was 10 h.

[0065] Example 3: S1 is made of 7A52 aluminum alloy rod with a diameter of 20 mm and 6061 aluminum alloy as the base material.

[0066] S2, using AFSD for layer-by-layer deposition, with deposition parameters meeting the following requirements: axial load of 40 kN, spindle speed of 800 rpm, travel speed of 10 mm / min, feed rate of 3 mm / min, and agitator shoulder diameter of 65 mm. A parallel three-layer overlapping path was used for layer-by-layer deposition to obtain a 7A52 aluminum alloy additive body. S3, solution treatment of 7A52 aluminum alloy additive body: heat to 490℃ and hold for 1 h, then water quench; The solution-treated additive body was subjected to a two-stage aging treatment: the first stage aging temperature was 115℃ and the holding time was 4 h; followed by the second stage aging temperature was 120℃ and the holding time was 20 h.

[0067] Comparative Example 1: S1 is made of 7A52 aluminum alloy rod with a diameter of 20 mm and 6061 aluminum alloy as the base material.

[0068] S2 uses AFSD for layer-by-layer deposition, with the following deposition parameters: axial load of 10 kN, spindle speed of 1300 rpm, travel speed of 50 mm / min, feed rate of 9 mm / min, and agitator shoulder diameter of 70 mm. A parallel three-layer overlapping path is used for layer-by-layer deposition to obtain the intermediate additive body. S3, Solution treatment of intermediate additive body: Heat to 470℃ and hold for 2 h, then water quench; The solution-treated additive body was subjected to a two-stage aging treatment: the first-stage aging temperature was 105℃ and the holding time was 8 h; followed by the second-stage aging temperature was 130℃ and the holding time was 14 h, to obtain the finished alloy.

[0069] Figure 11The tensile properties of the alloy obtained in Comparative Example 1 are shown in the diagram. The additive body exhibits low strength and ductility, with UTS and EL of 370 MPa and 13%, respectively. Comparative Example 1 employs a lower axial load combined with a higher rotational speed, faster travel, and larger feed rate, resulting in insufficient effective stirring and compaction of the material in the lap area. This leads to the presence of residual oxide films / weak bonding zones and micropores. These defects cannot be eliminated through subsequent solid solution and aging, becoming early cracking sources during the tensile process, causing a simultaneous decrease in strength and ductility. Furthermore, an unfavorable thermo-mechanical history exacerbates the inhomogeneity of the microstructure and solute distribution in the thickness direction, reducing the effect of two-stage aging on the uniform nucleation and controlled growth of the GP zone / η′, thereby weakening the aging strengthening efficiency. Therefore, Comparative Example 1 still exhibits lower UTS and EL under the same heat treatment regime. The staged control of this invention achieves controllable precipitate size and distribution within the deposition window that ensures dense forming and metallurgical bonding, thereby obtaining a stable strength-ductility match.

[0070] Comparative Example 2: S1 is made of 7A52 aluminum alloy rod with a diameter of 20 mm and 6061 aluminum alloy as the base material.

[0071] S2 uses AFSD for layer-by-layer deposition, with the following deposition parameters: axial load of 30 kN, spindle speed of 1000 rpm, travel speed of 20 mm / min, feed rate of 5 mm / min, and agitator shoulder diameter of 60 mm. A parallel three-layer overlapping path is used for layer-by-layer deposition to obtain the intermediate additive body. S3, Solution treatment of intermediate additive body: Heat to 470℃ and hold for 2 h, then water quench; The solution-treated additive body was subjected to a single-stage aging treatment: the aging temperature was 105℃ and the holding time was 22 h to obtain the finished alloy.

[0072] Figure 12The tensile curves and tensile properties of the finished alloy obtained in Comparative Example 2 are shown. The additive body exhibits low strength-plasticity, with UTS and EL of 351 MPa and 19%, respectively. Comparative Example 2 and Example 1 are identical in deposition parameters and solution treatment, differing only in the aging path. Comparative Example 2 employs single-stage aging (105℃ × 22 h). At the same temperature, the nucleation and growth processes are coupled, making it difficult to simultaneously achieve synergy between high-density nucleation and controlled growth / transformation. Firstly, although intragranular GP regions can form, the transformation to the η′ phase and the driving force for growth are insufficient and spatially uneven, leading to a decrease in the effective volume fraction and number density of η′ and a widening of its size distribution. Consequently, precipitation strengthening efficiency decreases, as evidenced by a significantly lower UTS than in Example 1. Secondly, prolonged single-stage aging more readily promotes preferential precipitation at grain boundaries and the formation / maintenance of PFZ, resulting in insufficient local strengthening near the grain boundaries and making them sensitive areas for strain concentration and microcrack initiation, making it difficult to obtain stable strength-plasticity synergy. In contrast, Example 1 pre-places GP regions extensively within the crystal as nucleation cores through first-stage aging, and second-stage aging promotes controlled growth of the cores and the formation of fine η′ phases (approximately 20 nm) with a more uniform distribution, while mitigating continuous precipitation at grain boundaries and PFZ problems, thereby achieving a better strength-plasticity match.

[0073] The above description is merely a preferred embodiment of the present invention and is not intended to limit the technical solution of the present invention in any way. Those skilled in the art should understand that, without departing from the spirit and principles of the present invention, the technical solution can be modified and replaced in several simple ways, and these modifications and replacements are all within the scope of protection covered by the claims.

Claims

1. A method for preparing high-strength and ductile 7xxx series aluminum alloys, characterized in that, include: S1, Clean the surface of the 7A52 aluminum alloy rod and aluminum alloy substrate; S2, using the cleaned 7A52 aluminum alloy rod as the deposition material, the 7A52 aluminum alloy additive body is obtained by layer-by-layer deposition on the cleaned aluminum alloy substrate using the friction stir deposition additive method. S3, the 7A52 aluminum alloy additive body is subjected to solution treatment and two-stage aging treatment in sequence to obtain the finished high-strength and ductile 7A52 aluminum alloy.

2. The method for preparing high-strength and ductile 7xxx series aluminum alloys according to claim 1, characterized in that, In S1, the chemical composition of 7A52 aluminum alloy bars by mass percentage is: Zn 4.0%~5.0%, Mg 1.8%~2.5%, Cu 0.1%~0.4%, Mn 0.2%~0.5%, Cr 0.1%~0.3%, Al 91.0%~93%, with the remainder being impurities.

3. The method for preparing high-strength and ductile 7xxx series aluminum alloys according to claim 1, characterized in that, In S2, the parameters for layer-by-layer deposition are: axial load of 20–40 kN, spindle speed of 800–1200 rpm, travel speed of 10–40 mm / min, feed speed of 3–8 mm / min, and shoulder diameter of the stirring head of 55–65 mm.

4. The method for preparing high-strength and ductile 7xxx series aluminum alloys according to claim 3, characterized in that, In S2, the deposited microstructure of the 7A52 aluminum alloy additive is a fine equiaxed crystal structure with an average grain size of ≤7 μm.

5. The method for preparing high-strength and ductile 7xxx series aluminum alloys according to claim 1, characterized in that, In S3, during solution treatment, the 7A52 aluminum alloy additive body is heated to 450-490℃ and held for 1-3 hours, and then quenched in water at 10-40℃ to obtain the intermediate aluminum alloy.

6. The method for preparing high-strength and ductile 7xxx series aluminum alloys according to claim 5, characterized in that, The intermediate aluminum alloy retains Fe and Mn enriched dispersed phases, including Al6(Fe, Mn) phase, with a content of 2.31%~3.36%.

7. The method for preparing high-strength and ductile 7xxx series aluminum alloys according to claim 1, characterized in that, In S3, the two-stage aging treatment involves sequentially performing a first-stage aging treatment and a second-stage aging treatment on the solution-treated 7A52 aluminum alloy additive body. During the first-stage aging treatment, the temperature is 90–115℃ and the holding time is 4–10 h to obtain a crude aluminum alloy. During the secondary aging treatment, the temperature is 120-150 ℃ and the holding time is 10-20 h to obtain a finished high-strength and ductile 7A52 aluminum alloy with η′ precipitates.

8. The method for preparing high-strength and ductile 7xxx series aluminum alloys according to claim 7, characterized in that, In S3, the equivalent size of the η′ precipitate is 15–30 nm.

9. The method for preparing high-strength and ductile 7xxx series aluminum alloys according to claim 7, characterized in that, Finished high-strength and ductile 7A52 aluminum alloy Al 18 The Mg3Cr2 phase contains nanotwin structures.

10. A high-strength, high-ductility aluminum alloy obtained by the preparation method of high-strength, high-ductility 7xxx series aluminum alloy according to any one of claims 1 to 9, characterized in that, The difference in ultimate tensile strength of this high-strength and ductile aluminum alloy between different deposited layers along the deposition height direction is ≤30 MPa, and the difference in elongation is ≤6%.

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