3.2.3. Microstructures in the Interior of Grains

Figure 8 shows the TEM bright-field images and high-resolution TEM (HRTEM) images of intragranular precipitates in the T6-treated matrix. In the 6061 alloy, all of the images were acquired in <001>Al directions, there are three types of precipitates observed: L phase, β" phase, and Q phase. The overview of known precipitates in the Al-Mg-Si alloy is listed in Table 2 in Ref. [3]. These precipitates were distinguished from morphology (Figure 8a) or with the help of HRTEM images and the corresponding Fast Fourier Filtering transform (FFT) patterns (Figure 8b–d). The needle-like precipitates were identified as the β" phase (Figure 8d), the lath-like precipitates with the cross-section elongated along <510>Al were the Q phase (Figure 8b), or with the cross-section elongated along <001>Al were the L phase (Figure 8c). The L phase is a precursor of Q phase. The addition of Cu has suppressed the formation of β" by the formation of Cu-containing precipitates (L and Q ), which result in the coexistence of lath-like and needle-like precipitates. In 7A52 alloy, all of the images were taken along <112>Al orientations, where a large number of GPII zones can be observed in Figure 8d. Two crystallographic orientations, <110>Al and <112>Al, are most suitable to identify disc-like precipitates in Al-Zn-Mg alloys, but along <112>Al more structural details can be revealed for the η -phase [38,39]. As shown in Figure 8d, the disc-like GPII zones are fully coherent with the Al matrix, composed of Zn-rich layers on the Al-{111} atomic planes. The profiles of GPII zones are presented by stress field contrast, resulting from atomic size difference between Zn, Mg, and Al.

**Figure 8.** Transmission electron microscope (TEM) images of precipitates: (**a**) bright-field TEM images of the T6-treated 6061 matrix along <001>Al zone axis; (**b**–**d**) high-resolution transmission electron microscopy (HRTEM) images of different types of precipitates in (a), they are Q´ phase (**b**), L phase (**c**), β" phase (**d**), respectively; (**e**) bright-field TEM images of the T6-treated 6061 matrix along <112>Al zone axis; and, (**f**) HRTEM image of GPII zones in (**c**). The inserts are the corresponding Fast Fourier Filtering transform (FFT) patterns.

Figure 9 shows the intragranular microstructure of WNZ in the weld joint of dissimilar T6-treated 6061 and 7A52. The corresponding hardness distribution across the weld joint is shown as the blue one in Figure 2. Figure 9a,b were taken from the area at the central line, where no precipitate is observed. Obviously, the pre-existing precipitates were resolved into the matrix during stirring, resulting in the lowest hardness being equal to the hardness of the solid solution state. The slight increase in hardness, from the 6061 side to the 7A52 side, might be due to solution strengthening increasing that results from alloying elements redistribution occurring during stirring. Over the lowest hardness on the 6061 side, the hardness increase is due to L precipitates forming, as shown in Figure 9c,d.

**Figure 9.** TEM bright-field images of the weld nugget zone (WNZ) along different zone axis: (**a**) on the 7A52 side along <001>Al zone axis; (**b**) on the 7A52 side along <112>Al zone axis; (**c**) on the 6061 side with the beam parallel to the <001>Al zone axis; (**d**) High-resolution TEM images and corresponding FFT patterns of L precipitate in (**c**) with the beam parallel to the <001>Al zone axis.

Figure 10 shows the microstructures after post-aging in the area at the central line in WNZ, where the hardness is in between that of the 7A52 and 6061 matrix, see Figure 4b. A large number of precipitates were observed along <001>Al and <112>Al orientations, as shown in Figure 10a,c, respectively. In Figure 10a, there are two types of precipitates with different size, the smaller with the size of about 1 nm are L phase, being identified from the HTTEM image in Figure 10b, the larger with the size of about 10 nm is difficult to identify from the <001>Al orientation. Projected along <112>Al (Figure 10c), the larger precipitates are also observed with distinct structural features, as shown in Figure 10d, they are identified as the η phase. Additionally, a few of stable η phase are also observed here. It must be noted that the coexistence of L phase and η phase in the same grain, which reveals the grain with the composition of Al-Zn-Mg-Si-Cu, it is the result of the fully mix of 7A52 and 6061 alloys by means of stirring. The above in agreement with the result revealed by the line scan in Figure 6.

**Figure 10.** TEM bright-field images of the same area in WNZ after post-aging as the same as Figure 4b along different zone axis: (**a**) along <001>Al zone axis; (**c**) along <112>Al zone axis. High-resolution TEM images and corresponding FFT patterns of different precipitates: (**b**) L precipitate in (**a**) along <001>Al zone axis; (**d**) η precipitate in (**c**) along <112>Al zone axis.

#### **4. Discussion**

The T6-treated 7A52 and 6061 alloys were joined by FSW, a region with the lowest hardness (Figure 2), the same as a solid solution, greatly weakened the weld joint. The lowest hardness results from precipitate decomposition (Figure 9a,b) during FSW under the role of heat inputs and dislocation moving. The dissolution of precipitates in WNZ is due to the temperature in FSW process is about ~425–480 ◦C, which is high enough to cause the dissolution of strengthening precipitates in WNZ [40]. An effective approach for improving the hardness of this region is to make precipitates form again. The local heat treatment would be difficult to employ, because this region is very narrow and the heat inputs will soften the adjacent area via precipitates coarsening. Thus, this study employed the integral heat treatment. For cost saving and improving elements diffusion, after solid solution treatment, the alloys were directly welded by FSW, and the weld joints were then aged. The displayed results revealed that the hardness of the WNZ is much higher than the matrix, as shown in Figure 3, which is satisfactory for the weld. However, the matrixes were significantly softened, because over-aging occurred during pre-heating at about 150 ◦C before the weld. Hence, the weld joints should be solid solution treated again.

When considering the different age temperature for 7A52 and 6061 alloys, the former is approximately 120 ◦C, and the latter is about 180 ◦C. Thus, there are three approaches for aging after post-solid-solution: (1) the welding structure was aged at 120 ◦C for 24 h and then 180 ◦C for 30 min.; (2) both sides of the weld joint were respectively aged: the 7A52 side was aged at 120 ◦C for 24 h, and then the other side 6061 was aged at 180 ◦C for 30 min.; and, (3) both sides of the weld joint were aged at 120 ◦C for 24 h, and then the 6061 side was further aged at 180 ◦C for 30 min. The ideal weld joints were obtained via the approach (1) the hardness change across the weld is described as in Figure 4b, where the 6061 matrix was strengthened to be similar to the T6-state, the 7A52 matrix was over-aged, and the hardness slowly changed in WNZ from 7A52 side to 6061 side. Composition analysis (Figures 5 and 6) and precipitates identification reveal the WNZ with the composition of Al-Zn-Mg-Si-Cu. During the first-stage aging at 120 ◦C, η-MgZn2-series precipitates form, and most of Zn and Mg moved from solid solution to form precipitates. The decrease of Mg in the matrix of WNZ inhibited the nucleation of β" phase, thus, after the second-stage aging, none of the β" precipitates were detected (Figure 10), unlike T6-treated 6061 alloy (Figure 8). In addition, Cu suppresses the formation of β" by the formation of Cu-containing L precipitates [3]. The hardness of WNZ is lower than that of 7A52 matrix and higher than that of 6061 matrix; it is dependent upon the mixed composition by 7A52 and 6061 alloys. In other words, the hardness is contributed by η and L precipitates. The number density of η precipitates is higher, the hardness of WNZ is closer to that of 7A52 matrix. Otherwise, the number density of L precipitates is higher, the hardness of WNZ is closer to that of 6061 matrix. Obviously, precipitates influence the hardness of the WNZ more.

The tensile properties of the weld that correspond to Figure 4 are shown in Table 2. After the post-weld heat treatment, the ultimate tensile strength of the weld significantly increased, in comparison to that of the weld without post-weld heat treatment; however, the elongations of all samples are very poor. The fracture topographies in Figure 11 shows a large number of coarse second phases (arrowed in Figure 11b,d) within dimples. In Figure 5, an amount of second phase is also observed on the grain boundary in WNZ. Obviously, these second phases formed during FSW, not being caused by the post-weld heat treatment. The micro-hardness is more dependent upon precipitates, while the tensile properties are usually affected by more complicated factors, such as defects, interfaces, the second phases, and so on.


**Table 2.** Tensile properties of the weld.

**Figure 11.** Fractured surface of the joint failed in the WNZ (SEM): (**a**) and (**b**) the sample of (SS + Welding); (**c**) and (**d**) the sample of (SS + welding + aging at 120 ◦C for 24 h + the second aging at 180 ◦C for 30 min.). (**b**) and (**d**) are the enlarged image of the boxed area in (**a**) and (**c**), respectively.
