Effects of Friction Stir Process and Stabilization Heat Treatment on Tensile Characteristics and Punch-Shear Properties of AZ61 Alloy
2017 Volume 58 Issue 1 Pages 6-10
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2017 Volume 58 Issue 1 Pages 6-10
In this study, the annealing AZ61 magnesium alloy (AO) was studied on friction stir process (AF). The AF specimen revealed that the non-uniform microstructure consists of fine and coarse grains in the stir zone. The FSPed specimens have shown improvement on microstructure and the particles are of matrix distributed over friction stir zone by stabilization heat treatment (AH). The specimens acquired a unique texture twins and basal slip system. The AF and AH specimens possess lower yield stress via friction stir processes and the uniform microstructure of the AH specimen increases the tensile ductility at tensile temperature of 100℃. On the other hand, the AH specimen has the amount of second phase particles and possesses great smooth surface on punch-shear test at 100℃.
Magnesium alloys have been considered as a candidate for aerospace industries as structural materials due to that they are the lightest of all metallic structural materials and possessing high strength-to-weight ratio1,2). Magnesium alloy has a poor formability at room temperature because of lack of the active slip from its hexagonal close-packed (hcp) crystal structure.
Friction stir weld welding (FSW) possesses more advantages than other traditional welding processes because it uses a solid-state joint process. After friction stir welding, structure is divided into stirred Zone (SZ), thermo-mechanically Affected Zone (TMAZ), heat Affected Zone (HAZ), and base Metal (BM), respectively3). In addition, friction stir processed is adopted as a material modification approach and is the method to refine the grain size. Our previous reports have investigated the tensile ductility of magnesium alloy up to 300℃, reported that the ductility can be improved at 100℃ due to the activity of non-basal slip, diversified twinning behavior and strain recovery effect4,5). It should be noted that the crystallographic orientation will resulting the decrease of yield stress of friction stir processed (FSPed) specimens. Compared to those hot extrusion specimens, the FSPed specimen can be acquired a unique strain hardening behavior on deformation process and pertains to the improvement of tensile ductility and so on. The FSPed specimens demonstrate a lower work hardening rate in the initial tensile deformation stage and maintain a prolonged strain hardening region before reaching the ultimate tensile stress when the specimens are tested at a boil water temperature of 100℃6). However, a significant difference in the strain hardening behavior between the hot-extrusion and the FSPed specimens resulted from the differences in the microstructure feature. Therefore magnesium alloy shows feasible for great working due to friction stir process.
In this research, the contribution of texture feature and twinning factor on the variation of tensile ductility will be the main focus of examination at room temperature (RT) and 100℃. Another aim explores the effect of some stabilization heat-treatment on the tensile ductility of hot-extrusion and friction stir processed AZ61 at RT and 100℃. Furthermore, the workability of magnesium alloys was investigated in the punch-shear process at room temperature and 100℃. The deformation behavior of punch-shear specimen was recorded in this study. The result of workability assessment has the value of the industrial application with magnesium alloy.
In this research, AZ61 Mg alloy has the chemical composition of 6.3 mass% Al, 0.48 mass% Zn, 0.37 mass% Mn. The 150 mm hot-extrusion billet was prepared by hot-top continuous casting, and subsequently extruded to plate at 350℃. The as-extruded plate was then subject to anneal at 345℃ for 12 hours (AO). Some AO specimens were machined to 100 mm (length) × 3 mm (thickness) × 25 mm (width) as the substrates for friction stir processed which is identified into. The processed direction (PD) was parallel to the extrusion direction (ED), with stirring tool which was at 1.5˚ from the process direction under a constant tool rotation speed (1241 rpm), the downward push force was controlled at 14.7 MPa and 60 mm/min traveling speed. The material of the backing plate is cast iron and the pin of the stirring tool had a diameter of 7 mm and depth of 1.5 mm, 16 mm in shoulder diameter. In addition, The AF specimens were subjected to stabilization heat-treatment at 200℃ for 1 hour. The stabilization heat treatment samples were designated as AH. The electron back scattered diffraction (EBSD) mappings were studied using scanning electron microscope (SEM, Zeiss Supra 55) equipped with EBSD that shows the IPF map and inverse pole figure. The scan step size is 0.5 μm and the accelerating voltage is 20 kV. The specimens obtained depth 1 mm from the top surface of the AO, AF and AH plates.
In this study, the length of the gauging section for this tensile specimen was 10 × 3.5 × 1.3 mm (Fig. 1(a)) and the tensile direction was parallel to the PD/ED. It should be noted that the gauge length portion of the FSPed specimen was completely located within the stirring zone. Uniaxial tensile test was carried out at constant initial strain of 1.67 × 10−3 s−1 at room temperature and 100℃. On the other hand, the location of the punch-shear test was observed in the middle of the FSPed specimens and parallel to the PD of the AF and AH specimens (Fig. 1(b)). The thickness of the specimen for punch-shear test is 1.5 mm. Figure 1(c) shows punch-shear equipment and the punch strain rate is 167 s−1. The punch-shear fracture surface of the specimens was observed that workability of magnesium alloys were probed at RT and 100℃.
Schematic illustration of the dimensions: (a) dimensions of the tensile specimens, (b) $ \overline{\rm AB} $ punching-shear surface, punch-shear section in the center of stir zone, (c) punch-shear equipment.
The microstructure of AO is shown in Fig. 2, which the grains have the average size of 13 mm. On the other hand, the macroscopic image of the ND observes the metal flow during the friction stir process (Fig. 2(b)). The non-uniform microstructure can be observed on the PD and normal plane (ND) as shown in Fig. 2(c), (d). The non-uniform microstructure was consist of fine grains (indicated with arrow in Fig. 2(b)), coarse grains, and the grain size, as of the fine grains were smaller than that of AO. This phenomenon is considered that the non-uniform heat remaining just after the FSPed and the fine grains of the dynamic recrystallization remained in the AF specimen. However, it is indicated that the macrostructure of metal flow is similar to the rotation trace of the stir pin because the lower heat input is introduced into the AZ61 specimen during friction stir process7). The microstructure of the AF specimen could be improved during stabilization heat-treatment, but the microstructure of ND also reveals a few inhomogeneous structures in the center of stir zone (Fig. 2(e), (f)). Comparing Fig. 3(a) with Fig. 3(b), showing that the friction stir processed causes the dissolution of γ-Mg17Al12, only few γ-Mg17Al12 and Al-Mn phase form in AF specimen. The EDS analysis of the particles in the AH specimens, the particles are γ-Mg17Al12 and Al-Mn phase in grain boundaries after stabilization heat-treatment as shown in Fig. 3(c).
Optical microstructures: (a) ED plane of the AO specimen, (b) the macroscopic image of the ND plane of the AF specimen, (c) PD plane of the AF specimen in the center of stir zone, (d) microstructures of inhomogeneous structure in ND plane of the AF specimen, (e) PD plane of the AH specimen in the center of stir zone, (f) ND plane of the AH specimen.
SEM micrographs: (a) γ-Mg17Al12 of the AO specimen(arrow), (b) the friction stir processed causes the dissolution of γ-Mg17Al12, (c) γ-Mg17Al12 and Al-Mn phase on grain boundary by stabilization heat treatment.
In addition, X-ray diffraction pattern (Cu target, 30 kV, 20 mA, 3˚/min) shows the (0002) direction is almost perpendicular to the ND in the AO (Fig. 4). In the friction stir process,8) the basal plane (0002) can be studied to be a column surface roughly surrounded the tool pin surface in the friction stir processed zone. The basal plane (0002) is roughly perpendicular to the PD in center regions. The IPF map and inverse pole figure with ED/PD for the AO、AF and AH specimen is shown in Fig. 5. A concentration of poles roughly near the (0002) direction indicates a strong basal texture as shown in Fig. 5(b). Yuan et al.9) obtained that the variation depth of the FSPed specimen shows different tilting angle of basal plane. It is reasonable to suggest that dimension of the tool pin causes basal plane (0002) cannot be completely parallel to PD.
X-ray diffraction patterns of the AO specimen, indicates high $(10 \bar{1} 0)$ and $(10 \bar{1} 1)$ intensities of the TD, (0002) perpendicular to the ND.
Microstructural features of materials, IPF map and inverse pole figure showing the orientation of the ED/PD: (a) AO specimen, (b) AF specimen, (c) AH specimen.
The results of tensile strength are shown in Fig. 6 (a), (b). Ultimate tensile stress (UTS) and yield stress (YS) of the AO specimens are higher than AF and AH at RT but the AO specimens possess lower tensile strength at a higher temperature. The previous research has mentioned that the orientation strengthening efficiency of Mg alloys is higher than the size comparison of the grain10). It is well known that the twinning system of $ \{10 \bar{1} 2 \} \langle 10 \bar{1} 1 \rangle $ only operate as c-axial tension. The basal planes of the AO specimen are almost parallel to the tensile direction, which is not favorable for the operation of basal slip system and $ \{ 10 \bar{1} 2 \} \langle 10 \bar{1} 1 \rangle $ tensile twinning11). The result of the AO specimen has a high flow stress due to the activity of non-basal slip systems and compression twins. The texture of the AF and AH specimen are favorable for the activity of basal slip system. The basal plane of some grains and tensile direction has a near zero included angle that is also favorable for tension twins. Figure 7 shows that the AO specimen has fewer twinning feature than others and it has the highest yield stress and the AF and AH specimens have similar yield stress for the domination of texture factor.
Tensile properties of AO, AF and AH: (a) YS, (b) UTS, (c) UE, (d) TE.
Subsurface of tensile specimens deformed to fracture: (a) AO specimen at RT, (b) AF specimen at RT, (c) AH specimen at RT, (d) AO specimen at 100℃, (e) AF specimen at 100℃, (f) AH specimen at 100℃.
As Fig. 2 (c), (d) shows, the AF specimen is beneficial to increase the ductility and the AH specimen is slightly decreasing at RT. In our recently study4), the orientation changing in the AZ31 FSPed samples can be inferred to the cause of the improved tensile ductility. The tensile ductility of the AF specimen is not as good as expected at 100℃ and the heterogeneous microstructure causes the similarity between the tensile ductility for two kinds of temperature in the AF specimen. There is lower heat input and even more severe inhomogeneous layer structures under the higher advancing speed condition. The non-uniform microstructure results in the poorer ductility at elevated temperature7). The stabilization heat-treatment could effectively improve the inhomogeneous microstructure of the FSPed specimens and then increased the tensile ductility at 100℃. Furthermore, both of the AO and the AF specimens has an uniform microstructure and it is beneficial to increase tensile ductility at 100℃. The microstructure factor contributes to the tensile ductility behavior of magnesium alloy at 100℃
3.3 Punch-shear property at RT and 100℃Figure 8 and Fig. 9 show punch-shear subsurface of AO, AF and AH specimens. The AO and AF specimens have rough punch-shear subsurface at RT and 100℃ as shown in Fig. 8 (a), (b) and 8 (c), (d) for AO and AF, respectively. Figure 9 (a), (d) shows the AO specimen is obviously observed large cracks parallel to fracture, and the AF specimen is observed rupture zone of large region from magnifying power. The AH specimen possesses the smooth punch-shear subsurface and correspond to the punch-shear tool shown in Fig. 8 (e), (f). Figure 9 (e) indicates that the punch-shear subsurface of the AF has few and thin cracks parallel to fracture at RT from magnifying power. Figure 9 (f) indicates that punch-shear subsurface has complete smooth and no crack in the AF at 100℃.
Subsurface of punch-shear specimens: (a) AO specimen at RT, (b) AF specimen at RT, (c) AH specimen at RT, (d) AO specimen at 100℃, (e) AF specimen at 100℃, (f) AH specimen at 100℃.
Subsurface of punch-shear specimens: (a) AO specimen at RT, (b) AF specimen at RT, (c) AH specimen at RT, (d) AO specimen at 100℃, (e) AF specimen at 100℃, (f) AH specimen at 100℃.
Figure 10 shows SEM micrographs of punch-shear fracture surface. The AO specimen can find the large cracks in two kinds of temperature as shown in Fig. 10 (a), (b), which the crack growth phenomenon can be improved after friction stir process (Fig. 10 (c), (d)). Figure 10 (e), (f) shows that fracture surface possesses greater smooth and the second phase particles can be seen. It can be observed that the rough fracture surface is the location of subsurface cracks by making a comparison between fracture surface and subsurface.
SEM micrographs of punch-shear fracture surface: (a) AO specimen at RT, (b) AF specimen at RT, (c) AH specimen at RT, (d) AO specimen at 100℃, (e) AF specimen at 100℃, (f) AH specimen at 100℃.
The cracks of the punch-shear fracture surface of AF were greatly improved after friction stir process, which is related with the texture and the punch-shear properties. It is noted that the second phase particles stay around the flat fracture surface and the crack region is without the second phase particles. The amount of flat surfaces is increased with increasing the punch-shear strength by increasing the amount of alumina particles in the aluminum matrix12). It is confirmed that the flat surfaces is closely related with the texture and the second phase particles (Fig. 10(e), (f)). The results of this study are similar to the punching test of LAZ92113) and the different magnesium systems used in those two tests show smooth subsurface after stabilization heat-treatment. On the other hand, the AF specimens possesses smooth subsurface and flat fracture surface at 100℃ that working temperature is also a major factor. The stabilization heat-treatment may be a great process for magnesium alloys FSPed to punch-shear process.
The authors are grateful to The Instrument Center of National Cheng Kung University, MOST 103-2221-E-006-056-MY2 for the financial support.
