**3. Results and discussions**

## **3.1 Tensile test**

During friction stir welding, the maximum temperature with plastic deformation causes an alteration in the distribution of the precipitate present in the base materials. So, the strength and ductility of the joints are also altered by the heat and temperature distribution in the welding process [17, 18]. Moreover, the FSW joints produce better joint efficiency compared to fusion welding processes, but there is a considerable difference between the base metal and the weld metal strength values.

*Center Stir Zone Investigations of Dissimilar AA6082, AA2014 and AA7075 Welds DOI: http://dx.doi.org/10.5772/intechopen.102652*

**Table 4** and **Figure 3** is showing the tensile test properties of the dissimilar joints welded using the FSW parameters of 80 mm/min, 900 rpm with tool tilt angle of 2°. It showed the AA2014/AA7075 joint with high tensile strength of 312 MPa than AA6082/AA7075 with the low tensile strength of 203 MPa. The AA6082/AA2014 weld showed an intermediated strength of 218 MPa. However, the AA6082/AA2014 dissimilar joint had a weld efficiency of 67%, which was slightly higher than the AA6082/AA7075 dissimilar joint, which had a weld efficiency of 62%. The higher tensile efficiency was seen for AA2014/AA7075 joint. Further, the AA6082/AA2014 joint had showed a 13% elongation percentage, whereas the AA6082/7075 joint had only a 7.5% elongation percentage.

**Figure 4** is showing the tensile failure samples. Due to the increased heat along the advancing side, it was also noted that all of the weld failures were located near the advancing side of the dissimilar joint. In addition to tensile failure, the SEM analysis were revealed approximately 98% of fine and shallow dimples confirmed the ductile mode of fracture of the weld (see **Figure 5**).

#### **3.2 Macrostructure**

Mostly friction stir welded macrostructure reveal geometrical effect of the tool in the central stir region as the reason for the onion ring formation in the weld [19]. Similarly the forces around pin forming center zone [5, 20] and defects such as tunnel, groove, kissing bonds, lack of diffusion, and flush in heat input or stirring during the process [21–23].

Here, **Figure 6** is showing macrostructure of the dissimilar joints of AA7075, AA6082, and AA2024. It shows a clear stir zone with defect-free joints. The dissimilar joints clearly showed the two materials mixed up in the joint's central stir zone. It is also observed the clear stir zone (SZ), thermo-mechanically affected zone (TMAZ), heat affected zone (HAZ), and base metal region (BM) on both sides of the joint material. Here, the basin shape center stir zone could be seen for all the dissimilar AA6082/7075, AA6082/2014, and AA2014/AA7075 joints.

## **3.3 Microstructure**

**Figures 7** and **8** are showing the microstructure variations of advancing and retreating sides of the AA6082/7075 weld. Also **Figures 9** and **10** are showing the microstructure of advancing and retreating sides of the AA6082/2014 weld. **Figures 11** and **12** show the microstructure of the advancing and retreating sides of the AA2014/7075 weld. Here stir zone, TMAZ, HAZ, and base material zones were revealed on the advancing and retreating sides of the microstructures.


**Table 4.**

*Tensile test properties of the dissimilar friction stir welded samples.*

**Figure 3.** *Tensile test of dissimilar welds.*

**Figure 4.** *Tensile failured samples.*

*Center Stir Zone Investigations of Dissimilar AA6082, AA2014 and AA7075 Welds DOI: http://dx.doi.org/10.5772/intechopen.102652*

#### **Figure 5.**

*SEM fractographs of dissimilar welds. (a) AA6082&2014, (b) AA6082&7075, and (c) AA2014&7075.*

#### **Figure 7.**

*Microstructure of advancing side of the AA6082/7075 weld. (a) AA6082BM, (b) HAZ, (c) TMAZ, and (d) stir zone.*

#### **Figure 8.**

*Microstructure of retreating side of the AA6082/7075 weld. (a) AA7075BM, (b) HAZ, (c) TMAZ, and (d) stir zone.*

*Center Stir Zone Investigations of Dissimilar AA6082, AA2014 and AA7075 Welds DOI: http://dx.doi.org/10.5772/intechopen.102652*

**Figure 9.** *Microstructure of advancing side of the AA6082/2014 weld. (a) AA6082BM, (b) HAZ, (c) TMAZ, and (d) stir zone.*

Jata et al. [24] clarified that the evolution of the microstructure of the Al-Li alloy welding zone consisted of 9 μm of grains due to dynamically recrystallized grains. Su et al. [25] established that the different zones of the FSW AA7050-T651 alloy as dynamically recrystallized zone (DXZ), thermo-mechanically affected zone (TMAZ), and heat-affected zone (HAZ).

In the microstructures below, the base metals AA6082 showed α-Al-Si-Mg solid solutions with their intermetallics. Similarly, AA7075 and AA2014 showed Al-Cu and Al-Zn α-solid solutions with their intermetallics. It was due to dynamic recovery and recrystallization in stir zone caused by high-speed rotation of the pin from top to bottom of the stir zone, revealed fine grain structure. The TMAZ zone has experienced heat and metal flow to reveal aligned intermetallics with deformed grains along the pin rotation are visible. Then, the HAZ zone with similar grains as that of base materials, which was influenced by the heat changes of welding, leading to some microstructure changes of coarse intermetallics.

## **3.4 Metal flow in stir zone**

Mukhopadhyay [26] highlighted the XRD and EDAX analysis of the precipitate phase-type/shape, habitat plane, and solute content for AA7075, AA2014, and AA6082. It was showed that the AA7075 (Al-Zn-Mg-Cu) with (Mg2Zn)/Plate in {111} plane, AA2014 (Al-Cu-Mg-Si) with S (Al2CuMg)/lath and GPZ/rod in {210} plane, and AA6082 (Al-Mg-Si) with (Mg2Si)/rod in <100> direction. Carron [27] highlighted the high peak aged (Mg5Si3-needles) up to 40 × 40 × 350 Å precipitates formation, before

#### **Figure 10.** *Microstructure of retreating side of the AA6082/2014 weld. (a) AA2014BM, (b) HAZ, (c) TMAZ, and (d) stir zone.*

 precipitate. After that, 10–20 μm of over aged (Mg2Si-rods) precipitates will form. The common precipitates sequences of the base metals are as follows: for the AA6082 supersaturated solution solid solution (SS) → Guinier-Preston (GP) zone (spherical) → (Mg5Si3-needle) → (Mg2Si-rod) → , for AA7075 supersaturated solution solid solution → GP zones → η′ (MgZn2) → η (MgZn2), for the AA2014 supersaturated solution solid solution (SSSS) → GP (rod) → θ<sup>1</sup> (Al2CuMg-lath) → θ (Al2CuMg) [28].

Mironov et al. [29] investigated the formation of high angle grain boundaries (HABs) on the Ti-6AL-4V alloy using friction stir processing (FSP) in microstructural studies. These LABs/HABs were clearly defined using electron back scattered diffraction (EBSD) studies. It was also stated that the α-titanium which was at room temperature in the HCP structure has to slip along the HCP crystal's basal plane to produce more HABs than LABs with dynamic recovery and recrystallization. So, it was therefore understood assertion that the material shear or strain low to moderate to high during deformation was mostly dependent on how the plane slides along a particular slip plane to create a substructure of low angle boundaries (LAB)/high angle boundaries (HAB).

Similarly in friction stir welding also the high speed rotation of the pin in the center stir zone shears or strains the material low to moderate to high deformation was mostly dependent on how the plane slides along a particular slip plane to create a substructure of low angle boundaries (LAB)/high angle boundaries (HAB). Kalemba-Rec et al. [13] explained the misorientation angle laid in between 2 and 15° was classified as low angle grain boundary (LAGBs) and whereas the misorientation angle greater than 15° was classified as a high angle grain boundary (HAGBs). So the grain

*Center Stir Zone Investigations of Dissimilar AA6082, AA2014 and AA7075 Welds DOI: http://dx.doi.org/10.5772/intechopen.102652*

**Figure 11.** *Microstructure of advancing side of the AA2014/7075 weld. (a) AA2014BM, (b) HAZ, (c) TMAZ, and (d) stir zone.*

size and its orientation variations could be clearly identified by electron back scattered diffraction (EBSD) scan images. Here, EBSD scan on zones such as SZ, TMAZ, HAZ, and base metal zones across the weldment gives the grain size variation across the AA6082/7075 and AA6082/2014 welds.

**Figure 13(a)** and **(c)** were showing the EBSD maps of the as-received materials of the base material AA6082-T6 with 53.3% of LABs and 46.7% of HAGs and the AA7075-T651 with 21.8% of LABs and 78.2% of HAGs. After welding base materials were not affected by any microstructural changes, thereby these base materials were not affected by the mode of metal flow. Moreover, in **Figure 13(b)** and **(d)** showing 42.5% of LABs and 57.5% of HABs of misorientation variations were seen on the advancing AA6082WN and 43.2% of LABs and 56.8% of HABs on the AA7075WN. In addition to misorientation variations, **Figure 14(a)** and **(b)** by EBSD analysis were showed 10–50 μm of grains in AA7075BM and 4–3 μm of grains in AA6082BM. Also **Figure 14(c)** and **(d)** were revealed the grains of 4–10 μm of grains on the AA6082 stir zone and 4–10 μm of grains on the AA7075WN were seen. Thus, the high speed rotating pin modified the grains to 4–10 μm of fine recrystallized grains in the stir zone were revealed in AA6082/7075 weld.

Similarly, the EBSD image quality figure with grain boundary orientation map of the AA6082-T6 and AA2014-T6 base metal used in the present study was shown in **Figure 15(a)** and **(c)**. It revealed that the total fraction of LAGBs was 48.3% and HAGBs was 51.7% on the advancing side of the AA6082, whereas the total fraction of LAGBs was 29.8% and HAGBS was 70.2% on the retreating side of the AA2014 side. In **Figure 15(b)** and **(d)**, the center stir zone revealed that the total fraction of LAGBs was 42.9% and HAGBs was 57.1% on the advancing side of the AA6082, whereas the

**Figure 12.** *Microstructure of retreating side of the AA2014/7075 weld. (a) AA7075BM, (b) HAZ, (c) TMAZ, and (d) stir zone.*

**Figure 13.** *EBSD misorientation images on: (a) AA6082BM, (b) AA6082WN, (c) AA7075BM, and (d) AA7075WN.*

*Center Stir Zone Investigations of Dissimilar AA6082, AA2014 and AA7075 Welds DOI: http://dx.doi.org/10.5772/intechopen.102652*

**Figure 14.**

*Grain size images on: (a) AA6082BM, (b) AA7075BM, (c) AA6082WN, and (d) AA7075WN.*

**Figure 15.**

*EBSD inverse pole figure (IPF) maps: (a) AA6082BM, (b) AA2014BM, (c) AA6082WN, and (d) AA2014WN.*

total fraction of LAGBs was 45.9% and HAGBS was 54.1% on the retreating side of the AA2014 side. The **Figure 16(c)** and **(d)** also confirms the 4–10 μm of fine dynamically recrystallized grains in the stir zone in AA6082/2014 weld.

Here, in this research, the AA7075 base material would deform along the {111} habit plane, AA2014 slips along {210} planes and AA6082 slips along the <100> directions, respectively by dynamically recovery and recrystallization. Thus, more grain substructures within the grain would evolve as a result of this dynamic recrystallization, which is surrounded by the high-energy grain boundaries of the weld matrix. As a result, substantial low angle grain boundary substructures are relatively common in the stir zone.

A vast hardness variations could be seen across the FSW weld. Since it results in advancing and retreating side BM, HAZ, TMAZ, and WN depend not only on the material and its rolled condition but also on the plastic deformation, frictional temperature and pressure applied during stir welding. Thomas et al. [15] and Threadgill et al. [16] studied the microhardness profiles measured in the welded material frequently exhibit a W-shaped characteristic. Moreover, Giraud et al. [30] experimented on asymmetric "W" shaped hardness profile of all joints, but it became cyclic on the welding path on the similar alloys AA7020-T651 to AA6060-T6.

In **Figure 17(a)**-**(c)** illustrated the hardness variations across the AA6082/2014, AA6082/7075, and AA2014/AA7075 weld, respectively. The hardness across weld was taken at t/2 (t-thickness) middle layer of the weld. The base metal hardness was 103.9 HV0.3 on AA6082BM and 121.5 HV0.3 on AA7075BM. Whereas, in **Figure 17(b)** 88.6 HV0.3 of hardness was observed in the center SZ, which was lower hardness than the base metal regions. This decreased hardness was due to the dynamic recrystallization with frictional heat in that zone. Consequently, 74.3 HV0.3 of hardness was

**Figure 16.** *Grain size on: (a) AA6082BM, (b) AA2014BM, (c) AA6082WN, and (d) AA2014WN.*

*Center Stir Zone Investigations of Dissimilar AA6082, AA2014 and AA7075 Welds DOI: http://dx.doi.org/10.5772/intechopen.102652*

#### **Figure 17.**

*Hardness variations across the dissimilar joints. (a) AA6082&2014, (b) AA6082&7075, and (c) AA2014&7075.*

noted on the advancing side AA6082HAZ, due to heat, which coarsened the grains with its precipitates in that zone. Thus the low hardness led to the tensile failure on AA6082HAZ side.

Similarly, in **Figure 17(a)** the initial observation showed that the hardness of the AA6082-T6 base metal was comparatively higher hardness than that of AA2014-T6 base metal. The hardness of the SZ matrix of the dissimilar weldment was lesser than that of BMs. Irrespective of the advancing side and retreating side of the dissimilar weldment of HAZ showed relatively lesser hardness than that its weld and its respective base metal. It was mainly attributed to the heating effect during FSW and thereby the disappearance of GP zones and the formation of overaged precipitates.

In **Figure 17(c)** the AA7075 base metal region showed high hardness than AA2014 BM side. Also the center stir zone showed the increased hardness from AA2014 advancing side to AA7075 retreating side. Also the hardness drop near AA2014 HAZ of 92.5HV0.3 had showed weld failure also.

The significant hardness drops in the HAZ were conceded that the maximum peak temperature evolved in the HAZ of the dissimilar weldment was touched above the β 1 and θ<sup>1</sup> and thereby the partial reversion occurs [31]. The literature confirmed that the main strengthening precipitates for the AA6082 AA were Mg5Si6 which was more stable at temperatures lesser than 200°C [32].

Generally, the precipitates were more stable and exist in the unaffected base material matrix and were absent in the SZ matrix and HAZ of the weldment. It was mainly attributed to the evolution of higher peak temperature in the SZ matrix during the heating cycle of FSW and it caused the dissolved strengthening in the weld matrix.
