*4.3. Influence of Mg-Rich Phase Particles on Superplastic Tensile and Fracture Process*

At 500 ◦C and 1 × 103 s–1, the components of the precipitated phase were analyzed using X-ray diffraction (XRD) and energy dispersive spectrometry (EDS), as shown in Figure 11. Full recrystallization of the deformed fine-grained 5A70 alloy structure after cold rolling was obtained using heat treatment. Figure 11a shows the dispersive fine precipitates in the recrystallized structure of the rolling surface. In addition, Figure 11b shows the results of the particle determination for the superplastic fracture specimens. There was a transition between the metastable β-Al2Mg phase and the β-Al3Mg2 phase with a hexagonal structure in the dynamic stretching process, as shown in Figure 11c [65]. Figure 11b,d indicate that the Al6(MnFe) phase disperses after all the heat treatments [30,66], and the Mg5Si6 phase precipitated during the superplastic deformation. Similar dispersed phase particles were previously observed in the aluminum-magnesium-silicon alloys [67,68].

**Figure 11.** Formation and composition of the dispersed phase particle morphology at full recrystallization (**a**), and the EDS/X-ray diffraction results of Al6(Mn,Fe) phase (**b**), β phase (**c**), and Mg5Si6 phase (**d**).

Precipitation hardening, where small particles inhibit the movement of dislocations to strengthen aluminum alloys, was used to improve the mechanical strength of Al-Mg alloys. The composition and structure of the Mg5Si6 phase in aluminum alloys were determined, which occurred as precipitates, and were associated with a particularly strong increase in the mechanical strength. This was due to the magnesium content increasing to 5.72 wt.%, while the addition of Mn, Fe, Cr, and Si improved the nucleation of the second phase particles [69,70]. However, the dynamic recrystallization was intensely impacted by the increased temperatures, and the dispersed distribution of the precipitated phases effectively impeded the grain growth and promoted the equiaxial transformation of the fine-grained structure, as shown in Figure 10. At 550 ◦C and 1 × <sup>10</sup>−<sup>3</sup> <sup>s</sup>−1, the inhibiting effect on the grain growth during tensile of the dynamic recrystallization decreased significantly when the grain size was 21.16 μm, as shown in Figure 10d. The abnormal grain growth resulted in an increase of the true stress under strain hardening (Figure 4b). The pinning effect of the precipitate was the intrinsic mechanism for the change in the true stress-true strain during the superplastic deformation and had a positive effect on the cavity nucleation and growth.

At 500 ◦C and 1 × <sup>10</sup>−<sup>3</sup> <sup>s</sup>−1, when the accumulated applied stress reached a maximum during the superplastic tensile, it was found that sliding of the grain boundary under the shear stress caused dislocations to pin up at the head of the phase particles, as shown in Figure 12a, and the cross grain boundaries slid and climbed to form a sub-grain boundary, as shown in Figure 12b. The TEM results suggest that the dislocation density in the fine-grained structure of the 5A70 alloy during superplastic deformation was ~5 × <sup>10</sup>−<sup>14</sup> <sup>m</sup>2. In addition, the dislocations gradually moved towards and were absorbed by the sub-grain boundaries during superplastic deformation. When the pile-up stress, *σp*, exceeded the theoretical decohesion strength of the Al-matrix/second phase particle interface boundary), the cavity began to nucleate [36]. At 500 ◦C and 1 × <sup>10</sup>−<sup>3</sup> <sup>s</sup><sup>−</sup>1, the maximum applied stress (*σmax* = 3.75 MPa) is substituted for the typical product stress *σ<sup>p</sup>* = 7.87 MPa [71]. Since the plugging stress threshold was more than twice the applied stress during superplastic tensile, the Al-matrix and the strengthening phase particles were easily separated and promoted cavity nucleation.

**Figure 12.** Location of dislocations, precipitated particles (**a**), and grain boundaries (**b**) in TEM images at *<sup>T</sup>* = 500 ◦C and . *<sup>ε</sup>* = 1 <sup>×</sup> <sup>10</sup>−<sup>3</sup> <sup>s</sup>−<sup>1</sup> with *<sup>ε</sup>* = 0.65.

The cavity nucleation and growth mechanism of the superplastic tensile was the cavity growth of stress promoting the diffusion of the small cavity along the grain boundary, including the superplastic diffusion growth, and the plastic-controlled growth caused by the plastic deformation of the recrystallized grain around the cavity [72]. The pinning effect of the precipitated phase enhanced the cavity nucleation and the chemical potential between the forceful grain boundary atoms and the free surface of the cavity at 550 ◦C [73].

Cavity expansion mainly occurred during the growth of the diffusion and the superplastic diffusion because of the spreading of voids into the nearby cavity since the cavity radius was less than the grain size. In addition, plastic-controlled growth dominated the cavity interlinkage and coalescence process, which eventually led to superplastic fracturing [74]. The expansion of the cracks was the fundamental reason for the transient instability and superplastic fracture. The behavior of cavities nearby the fracture surface was studied, as shown in Figure 13, which illustrates the fracture morphology of the superplastic tensile specimens.

**Figure 13.** Fracture morphologies and cavity behaviors for the superplastic fraction surfaces of the 5A70 alloy at *<sup>T</sup>* = 400 ◦C, . *<sup>ε</sup>* = 1 <sup>×</sup> <sup>10</sup>−<sup>3</sup> <sup>s</sup>−<sup>1</sup> (**a**); *<sup>T</sup>* = 450 ◦C, . *<sup>ε</sup>* = 1 <sup>×</sup> <sup>10</sup>−<sup>3</sup> <sup>s</sup>−<sup>1</sup> (**b**); *<sup>T</sup>* = 500 ◦C, . *<sup>ε</sup>* = 5 <sup>×</sup> <sup>10</sup>−<sup>4</sup> <sup>s</sup>−<sup>1</sup> (**c**); and *<sup>T</sup>* = 550 ◦C, . *<sup>ε</sup>* = 5 <sup>×</sup> <sup>10</sup>−<sup>4</sup> <sup>s</sup>−<sup>1</sup> (**d**).

Figure 13a–d show that no significant necking occurred at the superplastic fracture, and cracks near the fracture location gradually spread outwards from the fracture along the accumulated deformation. Therefore, crack formation was the main reason for the ultimate fracture of superplastic tensile [75]. The pinning effect of the phase particles was strengthened to enhance the cavity nucleation and the chemical potential between the forceful grain boundary atoms and the free surface of the cavity. In a previous study, the results demonstrated a clear transition from diffusion growth to superplastic diffusion growth and plastic-controlled growth at a cavity radius larger than 1.52 and 13.90 μm [71]. Cavity growth mainly occurred at the stage of diffusion growth and superplastic diffusion growth due to the diffusion of the voids into adjacent cavities with a cavity radius smaller than the grain size. Plastic-controlled growth dominated the cavity interlinkage and coalescence process, which eventually led to superplastic fractures. Therefore, the irregular-shaped cavity with the accumulation of deformation strains induced wedge cracks, which propagated and converged under the shear stresses until a fracture occurred.

At the first stage, the crack in the specimen propagated as an opening mode crack that formed vortex structures (Figure 7). Second phase particles were formed along the cavity surface, which indicates a high mass transfer rate in the vortex porous structure. At the second stage, periodic transverse shear displacement occurred at the fracture surface (Figure 13b,c) due to a change of the crack type: the crack at this stage propagated as a sliding mode crack. As the crack reached the specimen edges, it rotated and propagated along the direction of *τmax* under plane stresses. At this stage, the fracture fractograph revealed traces of material rotations along the longitudinal shear in the form of plane cavities and discontinuities (Figure 13d) [76].

At 550 ◦C, the dissolution of the phase particles reduced the suppression of grain growth during dynamic recrystallization. However, the grain growth promoted the healing of the small cavities [71]. This clearly verifies that the cavity nucleation and cavity growth during superplastic tensile deformation promoted steady-state flows. Ultimately, the cracks destroyed the superplastic tensile stability and resulted in superplastic fracturing with no obvious necking.
