Investigation of 7075 aluminum alloy TIG welding joint using 7075 aluminum alloy wire before and after heat treatment

7075 aluminum alloy butt welding was carried out by TIG welding with self-made 7075 wire. The effects of different currents on the microstructure and mechanical properties of the joint before and after heat treatment were studied. Before heat treatment, the strength and elongation of the joint are 274.84 MPa and 6.3%, which are determined by the presence of a large number of second phase precipitates with different shapes. After heat treatment, the tensile strength and elongation are increased to 511.75 MPa and 13.76%. Since most of the second phase has been dissolved, the grain size becomes the decisive factor. The fine grain size and more sufficient recrystallization degree make the mechanical properties of the low-current joint better.


Introduction
7075 aluminum alloy belongs to Al-Zn-Mg-Cu series heat-treatable aluminum alloy in 7XXX aluminum alloy, which has the advantages of high specific strength, good machining performance and high fracture toughness. It is widely used in aerospace, military and national defence, rail transportation, shipbuilding and other fields [1][2][3][4]. 7075 aluminum alloy is a precipitated hardening alloy, due to its poor weldability, easy hardening cracking and over-aging in the heat-affected zone, it is difficult to weld by traditional welding methods such as TIG and MIG, which has become a bottleneck problem restricting its welding application [5][6][7][8][9][10][11][12][13][14].
7075 aluminum alloy is generally weld with 5XXX aluminum alloy welding wire by fusion welding, the maximum joint strength can only reach 60% of the base metal [12]. The softening of 7075 aluminum alloy welded joint is unavoidable, and the greater the welding heat input, the more serious the joint softening [15]. After welding with 5XXX aluminum alloy welding wire, the weld performance of the joint cannot be improved by heat treatment, which has been the crux of the current 7075 high-strength aluminum alloy fusion-welding problem. According to the principle of equal strength selection of welding materials, and because of 7xxx aluminum alloy can precipitate the second phase in a specific sequence to strengthen the α matrix [16][17][18][19], it can be seen that 7XXX aluminum alloy welding wire is very suitable for 7075 welding. Oropeza D et al used self-made titanium carbide reinforced 7075 welding wire to TIG weld 7075 aluminum alloy, and the joint strength after heat treatment could reach up to 84% of the base metal [20]. Zhang K et al used self-made 7075 wire and TIG welding to weld 5 mm thick 7075-T6 aluminum alloy, and then heat-treated the joint to obtain welding joint with the tensile strength of 424.5 MPa and elongation of 9.83% [21].
In this paper, the butt welding of 5 mm thick 7075 aluminum alloy was carried out by TIG welding using 7075 wire formed by drawing out with self-made spray forming billet. After welding, the joint was heat treated. The microstructure and properties of the joints under different current parameters were compared and analyzed to study the influence of current on the microstructure and properties of the joints before and after heat treatment, so as to evaluate the feasibility of using 7075 welding wire produced by the new process in this paper to improve the joint properties by heat treatment after welding of 7075 aluminum alloy.

Materials and methods
7075 high-strength aluminum alloy test plate with the size of 5 mm × 150 mm × 75 mm was selected as the base metal, with the tensile strength of 545 MPa, yield strength of 475 MPa and elongation of 8%. The self-made 1.2 mm 7075 welding wire formed by drawing out with self-made spray forming billet was used as welding material, as shown in figure 1(a). The chemical composition of base metal and welding wire is shown in table 1.
The welding machine used was TIG 4300i AC/DC inverter argon arc welding machine produced by ESAB Company, the shielding gas was 99.99% pure argon, with 60°groove, 2 mm dull edge, welding speed of 3 mm s −1 , wire feeding speed of 200 cm min −1 , duty cycle of 60%, AC connection, frequency of 120 Hz, arc length of 3.2 mm. When the current is less than 160 A, single pass penetration cannot be achieved, and when the current is more significant than 180 A, macro cracks appear, so the welding currents were set at 160 A, 165 A, 170 A, 175 A and 180 A, respectively, to get the welding joint. Joint form and weld pass appearance are shown in figures 1(b), (c). As shown from figure 1(c), the weld was well-formed without apparent defects. With the increase of current, the weld surface appeared slight collapse, and the collapse was more evident at 180 A.
After welding, the joint was heat treated: stress relieving (230°C, 8h)→solution (480°C, 1h)→aging(120°C, 24h), and the joint samples before and after heat treatment were respectively intercepted in the direction perpendicular to the weld pass. Tensile samples are shown in figure 1(d). Tensile tests at room temperature with loading speed of 1 mm min −1 , and 6 parallel samples were tested for each weldment. The Vickers hardness of the joint with a load of 200 g and the pressure was held for 12 s. The location of the hardness test is shown in figure 1(e).

Results and discussion
3.1. Microstructure of joint Figure 2 shows the Microstructure of the welding joint under different currents before and after heat treatment. The observation area of each welding joint sample can be divided into base metal (BM), fusion zone (FZ), heataffected zone (HAZ) and weld zone (WZ), as shown in figure 2(a). Figure 2(b) shows that the BM has a typical rolled structure, and the grain is banded or fibrous along the deformation direction. After heat treatment, the BM structure grows to a certain extent. Figures 2(c)-(g) show the microstructure morphology of the joint at the welding current of 160 A-180 A before and after heat treatment respectively. It can be seen that the WZ before heat treatment has relatively uniform equiaxed crystal structure, due to the release and outward heat conduction of latent heat during crystallization and more alloying elements in the molten pool promoting nucleation, local thermal equilibrium is quickly achieved in the molten pool to form equiaxed crystals. The FZ is located between the weld and the HAZ, which is the starting position of weld crystallization. Because the maximum degree of supercooling occurs in the FZ, and the separated impurity elements will promote more non-spontaneous nucleation, resulting in a smaller grain in the FZ than the WZ. Before heat treatment, due to local melting, the  microstructure is not uniform. At the position where the FZ is close to the HAZ, due to the proximity to the BM, the cooling rate is very fast, so that the grain along the heat dissipation direction growth time is short, and then the formation of fine grain. At the position where the FZ is close to the weld, columnar crystals grow towards the center of the weld due to the large temperature gradient in the weld and develop into shorter columnar crystals due to the fast-cooling rate. The side of HAZ near the weld is a phase change recrystallization zone, with refined equiaxed grains evenly distributed. The side of HAZ near the BM is an incomplete phase change recrystallization zone, where the grains grow significantly larger than the BM, but still maintain the rolling state characteristics, and some recrystallization structures appear. Temperature is the main factor affecting grain growth. The higher the temperature is, the faster the grain boundary migrates and the higher the grain growth rate is. The welding current directly determines the joint temperature. Figures 3(a)-(d) show the grain distribution of the HAZ at 160 A and 180 A before and after heat treatment. The average grain size measured by the transversal method (divide the random transversal by the number of grains the transversal passes through) is shown in figure 3(e). The average grain size in the HAZ is 22.84 μm when the current reaches 160 A, and 24.76 μm when the current increase to 180 A. With the process of solid solution heat treatment, the second phase particles gradually dissolve back into the matrix, and the impeding effect of the second phase particles on grain boundaries gradually decreases, the migration rate of grain boundaries becomes faster, and the growth rate of grains increases, resulting in the growth of grain size after heat treatment. After heat treatment, the grain size of HAZ under current 160 A and 180 A grow to 25.21 μm and 26.7 μm, respectively.
Combined with figures 2 and 3, comparing microstructure under different currents of the weld, it can be seen that the grain size increases significantly with the increase of current for two reasons: (1) The increase of current corresponds to the increase of heat input, which leads to the extension of the residence time of the weld in the high-temperature zone, resulting in the coarsening of grain structure; (2) The increase of heat input leads to more serious burning loss of alloying elements, resulting in the reduction of non-spontaneous nucleation impurities and coarsening of grains. By comparing the microstructure before and after heat treatment, it can be seen that: (1) After heat treatment, the grain structure of the joint grows obviously, but there is no over-burning phenomenon; (2) The structure is more uniform after heat treatment than that before heat treatment; (3) Before heat treatment, the black precipitated phase in each area of the joint is distributed on the grain boundary in mesh layer, and the grain boundary is obvious. After heat treatment, the black precipitated phase basically disappears, and the grain boundary is not obvious, except for some dispersed granular second phase particles; (4) After heat treatment, the grain size still keeps the characteristic that the grain size increases with the increase of current, that is, the influence of welding heat input before and after heat treatment on the grain size is hereditary. Figures 4(a)-(d) shows the recrystallization distribution of HAZ under different current before and after heat treatment. The figure shows the recrystallized grains in blue, while the yellow and red grains represent the sub crystalline grains and deformed grains, respectively. Combined with the recrystallization distribution trend diagram in figure 4(e), it can be seen that, in the process of welding and heat treatment, partial recrystallization occurs in the grain structure of the HAZ. Before heat treatment, the recrystallization coefficient of the HAZ is  72.6% at 160 A current and 77.8% at 180 A current. After heat treatment, the recrystallization coefficient of the HAZ at 160 A current is 90.7%, and that at 180 A is 82.8%. Before heat treatment, as the heat input under 180 A current is larger than that under 160 A current, the recrystallization in the HAZ is more sufficient under high current due to the influence of the weld thermal cycle. In the process of heat treatment, the recrystallization process continues. As the grain structure is finer under a small current, the more nucleation positions of recrystallization and the increase of grain boundary coefficient, resulting in significant energy storage at the interface, the recrystallization degree of the joint structure under small current has higher-level sufficiency after heat treatment. Therefore, under the same heat treatment process conditions, the smaller the welding current, the more conducive to recrystallization. Figure 5 is the XRD pattern of the joint at each current before and after heat treatment. The most substantial peaks, such as (111), (200) and (220), correspond to the typical reflection peak of Al matrix. η(MgZn 2 ) phase is the main strengthening phase of 7075 aluminum alloy, and its reflection peak can be observed in all samples before heat treatment. Due to the effect of the thermal welding cycle, S(CuMgAl 2 ) phase, T(AlZnMgCu) phase, and Al 13 Fe 4 phase also appear in the joint samples before heat treatment. Due to its low content, only weak reflection peaks can be observed. After heat treatment, the reflection peaks of η(MgZn 2 ) phase, S(CuMgAl 2 ) phase, and T(AlZnMgCu) phase disappeared. Except for aluminum, only weak reflection peaks of AlCu x phase and Al 13 Fe 4 phase are displayed. Figure 6 shows the SEM images of the weld joint in the WZ and HAZ, in which figures 6(a)-(e) correspond to the joint under 160 A-180 A before heat treatment respectively, and figures 6(f)-(j) correspond to the joint under 160 A-180 A after heat treatment respectively. The light-colored part in the figure is the precipitated phase. Figures 6(a)-(e) show that there are a lot of second phase structures in the weld before heat treatment, and these second phase particles are densely distributed on the grain boundary, even in a continuous distribution state. The left side of the HAZ image contains the FZ and some BM. Due to the serious segregation of alloy elements in the FZ, some second phase particles are coarser and distributed in blocks. By comparing the weld, FZ and HAZ, it can be seen that the precipitated phase before heat treatment has similar characteristics in morphology, type and distribution. They are mainly composed of a large nodal phase (phase A), the strip phase (phase B) densely distributed on the grain boundary, which mostly network layered structure, large massive phase (phase C) and small granular phase (phase D). The SEM images and EDS analysis of typical characteristic phases are shown in figures 7(a)-(d) respectively. It is very complex for eutectic phase with a large amount of alloying elements. Some elements would be dissolved in the second phase, meanwhile, segregation of alloying elements would be occurred when components were welded, and affected by welding thermal cycle, these second phases would grow, resulting in junction. Hence, it can be only deduced the phases or the main phase [22]. According to the analysis of XRD data, for phase A, it is a eutectic structure with fine lamellae, and the main alloy elements are Zn, Mg and Cu. As the atomic radii of Zn and Cu are similar, the atomic ratio of Zn: Mg: Cu can be considered as about 1:1:1 according to the EDS analysis results, and phase A is judged to be a non-equilibrium binary eutectic structure formed by T(AlZnMgCu) phase and α(Al) phase; For phase B, the atomic ratio of Zn: Mg: Cu is about 2:2:1, so phase B is judged to contain η(MgZn 2 ) and S(CuMgAl 2 ) phases; For phase C, the atomic ratio of Fe is about 14.1%, so the dominant phase should be Al 13 Fe 4 or Fe containing compounds; For phase D, the atomic ratio of Zn: Mg: Cu is about the same as that in phase A, so it is judged to be a non-equilibrium binary eutectic structure formed by S(CuMgAl 2 ) or T(AlZnMgCu) phase and α(Al). Phase B and D appear in large numbers in the WZ, FZ and HAZ, which is the main forms of a precipitated phase of the joint; Phase A mainly appears in the FZ and rarely in the weld, but it tends to increase with the increase of current because the element segregation in the FZ is serious and the heat input increases with the increase of current, resulting in the growth of the second phase; Phase C mainly appears in the FZ and HAZ and rarely in the weld as the Fe content of the welding wire is only 0.03%. Due to the influence of the thermal welding cycle, the second phase particles of different types are mostly connected together and distributed on grain boundaries after precipitating and growing up, which seriously affects the mechanical properties of the joint. Figures 6(f)-(j) show that after heat treatment, the second phase in each area of the joint is significantly reduced, the network phase disappears, most of the second phase has been basically dissolved in the matrix, and only a small amount of dispersed rectangular residual phase is left. This is because for Cu containing alloy, when the Zn/Mg ratio is larger than 1.5, the second phase in the alloy is easy to dissolve [23]. Typical residual phase morphology and EDS analysis are shown in figures 7(e)-(g), strip E phase, granular agglomerated F phase and massive G phase respectively. Combined with XRD analysis, it can be seen that E and G phase are mainly composed of Al 13 Fe 4 and AlCu x , with AlCu x as dominated, while Al 13 Fe 4 dominates F phase. The agglomerated F phase mostly appears in FZ and HAZ, but not in the welds. It can be seen that η(MgZn 2 ), S(CuMgAl 2 ), the non-equilibrium binary eutectic structure formed by S(CuMgAl 2 ) or T(AlZnMgCu) phase and α(Al) have been completely dissolved after heat treatment of the joint, and the insoluble Al 13 Fe 4 and AlCu x phases are dispersed in the matrix. Figure 8 shows the area fraction of the second phase under different currents before and after heat treatment according to the statistics of Image-Pro Plus (IPP) software. Under the same welding parameters, the area fraction of the second phase between WZ and HAZ is very different, so the average area fraction of each zone is taken. Before heat treatment, with the increase of welding heat and decrease of the cooling speed of the weld, the precipitated phases such as η and S phases grow up, the area fraction of the second phase of the joint increases with the increase of current, from 6.4% (160 A) to 8.8% (175 A). At 180 A, the obvious increase of crystallization defects occupies part of the area statistics, therefore, the area fraction of the second phase decreases at 180 A. After heat treatment, precipitated phases such as η and S phases are completely dissolved, and the area fraction of the second phase decreases substantially overall. Due to the low content of insoluble residual phases such as Al 13 Fe 4 and AlCu x , there is little difference in the area fraction of residual phase from 160 A to 175 A, with an average of 1.16%. At 180 A, due to the serious crystallization defects and element segregation, the area of insoluble residual phase is large, resulting in the residual phase area fraction of 1.9%, which is larger than the average after heat treatment.

Analysis of tensile properties and micro-hardness of joints
Tensile results of welding joints under different current parameters before and after heat treatment are shown in figure 9. As can be seen from figure 9(a), the tensile stress-strain curve of the joint before heat treatment has no   Figure 9(c) shows that the elongation of the joint is also greatly improved after heat treatment. Before heat treatment, there is no obvious proportional correspondence between the joint elongation and the current change value, the maximum is 6.3% (175 A) and the minimum is 5.35% (170 A); After heat treatment, the elongation decreases with the increase of current, from 16.9% (160 A) to 7.9% (180 A).
The Vickers micro-hardness distribution of the cross-sections of the joints before and after heat treatment is shown in figures 10(a), (b). Before heat treatment, the hardness of the joint showed a 'W' distribution, and the average hardness of the weld was 125.4 HV. Several reasons include burning segregation of alloying elements in the FZ, the heat-affected area being over-aged, the precipitation phase growing up (figure 6), the grains growing up (figure 2), etc. As a result, the hardness of the FZ decreases sharply, and the position close to the fusion line drops to a minimum of 83 HV. The obvious decrease in the hardness of the FZ and the HAZ causes the softening of the joint and affects the performance of the joint, tensile fracture positions are all in this region. After the heat treatment, the cross-section hardness distribution no longer has the typical fusion weld 'W' distribution characteristics. The overall hardness of each zone of the joint and the BM has been dramatically increased to a similar level, with an average hardness of 165 HV, which eliminates the softening of the joint caused by the decrease in the hardness of the HAZ. It should be noted that the hydrogen solubility of aluminum alloy in liquid/ solid-state has a vast difference, and hydrogen evolution in the weld is inevitable. After heat treatment, hydrogen evolution pores and solidification defects are magnified (figure 2), resulting in the hardness of each joint under different currents fluctuating within 20 HV and the hardness value is unstable; At 180 A, the hardness of joint is irregular and often appears steep drop, which is caused by many solidification defects such as shrinkage cavity, porosity and microscopic crack when the current is high (figure 6(e)).

Conclusions
(1) 7075 aluminum alloy butt welding was carried out by TIG welding with self-made 7075 wire, and the joint without defects was obtained. After heat treatment, the mechanical properties of the joint have been greatly improved. The maximum elongation reaches 16.9% and the tensile strength reaches 511.75 MPa, which is 94% of the tensile strength of the base metal.
(2) Before heat treatment, η(MgZn 2 ), S(CuMgAl 2 ), T(AlZnMgCu) and other precipitates distributed continuously at the grain boundary reduce the mechanical properties of the joint and weaken the influence of grain size, leading to the formation and area fraction of precipitated phase become the determining factors of the mechanical properties.
(3) After heat treatment, most of the precipitated phase dissolved into the matrix, and the grain size became the main factor affecting the mechanical properties of the joint. The fine grain size and more sufficient recrystallization degree make the mechanical properties of the low-current joint better.

Acknowledgments
This work is supported by the National Natural Science Foundation of China (No. 5210130372). The author also thanks Xijing Wang in Lanzhou University of Technology for discussing the technology of welding schemes with the author.

Data availability statement
All data that support the findings of this study are included within the article (and any supplementary files).