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Evolution of grain structure, micro-texture and second phase during porthole die extrusion of Al–Zn–Mg alloy
Materials Characterization 158 (2019) 109953
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Evolution of grain structure, micro-texture and second phase during porthole die extrusion of Al–Zn–Mg alloy
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Liang Chen, Gaojin Chen, Jianwei Tang, Guoqun Zhao∗, Cunsheng Zhang Key Laboratory for Liquid-Solid Structural Evolution and Processing of Materials (Ministry of Education), Shandong University, Jinan, Shandong, 250061, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: Al–Zn–Mg alloy Extrusion Grain structure Texture Second phase
The porthole die extrusion of high strength Al–Zn–Mg alloy was carried out at 783 K. The reserved materials inside portholes and welding chamber were taken out to investigate the dynamic evolution of grain structure, micro-texture and second phase along streamlines of welding zone and matrix zone. The material of welding zone experienced severe plastic deformation, while the flowing route of matrix zone was relatively smooth. Hence, the welding zone achieved near complete dynamic crystallization after splitting stage, and the grain size kept steady during the subsequent welding and extruding stage. The streamline of matrix zone mainly consisted of elongated grains with small amount of fine equiaxed grains, which indicated the occurrence of dynamic recovery and slight dynamic crystallization. The texture evolution of welding zone was much more complicated due to the combination of compression, shearing and tension stresses. The welding zone of the final extruded profile had the main textures of {112} < 111 > , {101} < 111 > , {101} < 001 > and {111} < 112 > orientations, while the matrix zone of the profile consisted of {112} < 111 > , {101} < 111 > and {101} < 211 > orientations. The extrusion temperature is sufficient for dissolving MgZn2 phase, while it is insufficient for the dissolution of Al23CuFe4. Instead, the coarse Al23CuFe4 was broke into small pieces due to the effects of mechanical deformation.
1. Introduction High strength Al–Zn–Mg alloys have been widely applied in the fields of aerospace, high speed train and automobile, due to their advantages of low density and high specific strength [1–4]. Hot extrusion is one of the main forming processes on Al alloys to produce extruded profiles with identical cross-section. In order to obtain Al profiles with high dimensional accuracy and excellent mechanical properties, both the material flowing behavior and microstructure evolution during hot extrusion should be well controlled. In the past decade, the researchers have attempted to enhance the flowing homogeneity through optimizing the die structure and process parameters [5–8]. However, the evolution of microstructure during hot extrusion has not been clarified. As is known, the grain morphology, texture feature and second phase evolve markedly during hot extrusion process. The occurrence of dynamic recovery (DRV) and dynamic recrystallization (DRX) can help to reduce the grain size, and it is beneficial for improving the mechanical properties of extruded profiles [9,10]. Moreover, the texture evolution such as the type, density, and distribution also plays an important role on the final mechanical properties. It has been proved that the anisotropy of extruded Al profiles can be obviously enhanced by
∗
controlling the texture during hot extrusion [11]. In case of second phases, their composition, morphology and distribution should be well examined. In general, fine second phase with uniform and dispersed distribution can effectively improve the microstructure, while coarse second phase has adverse effects [12,13]. Flat die is usually used in practical production to produce solid Al profiles with simple cross-section. The researchers have carried out lots of experimental works to study the microstructure characteristics of solid Al profiles. Bois-Brochu et al. [14] examined the microstructure, texture and static mechanical properties of extruded Al–Li alloy, and found that the variation of strength and anisotropy was strongly dependent on billet temperature. Xu et al. [15] proposed that the size of eutectic Si particles and primary intermetallic particles of Al–Si–Cu alloy was reduced after hot extrusion, and the profile extruded at 450 °C exhibited the best combination of strength and ductility. Li et al. [16] proposed that the asymmetric feeder chamber can modify the grain structure, texture component and second phase distribution of the extruded Al–Zn–Mg alloy, which led to the decreasing anisotropy of elongation. Wu et al. [17] conducted extrusion experiment using Al–Si–Mg alloys with various Si contents, and found that the DRX behavior during hot extrusion was strongly related to the number of large Si
Corresponding author. E-mail address: [email protected] (G. Zhao).
https://doi.org/10.1016/j.matchar.2019.109953 Received 27 June 2019; Received in revised form 14 August 2019; Accepted 1 October 2019 Available online 16 October 2019 1044-5803/ © 2019 Elsevier Inc. All rights reserved.
Materials Characterization 158 (2019) 109953
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amount of second phases can be observed from Fig. 1(b). EDS data implies that the second phases located at triangle region of grain boundary are Al23CuFe4, while those distributed inside the grains are MgZn2. In addition, there are also some second phases, which are characterized as Al2CuMg, discontinuously distribute along grain boundary. As shown in Fig. 1(c), the peak texture intensity is as low as 2.77, which means that the homogenized alloy has no obvious orientation preference. The extrusion setup and the main dimension of the bridge, upper die, and lower die are schematically drawn in Fig. 2. The bridge is separated from upper die in order to facilitate the removal of reserved material inside portholes and welding chamber after extrusion experiment. Moreover, for convenience of taking specimens for microstructure observation, a plate shaped profile was extruded out. The homogenized billet with the dimension of ϕ40 × 60 mm was machined. Then, the billet was firstly heated to 783 K together with the die, and held for 20 min. After that, the extrusion was started, and the ram velocity is set as 6 mm/min. The experiment was stopped when the stroke reached 30 mm, and the whole setup was quenched by cooling water. Finally, the remaining materials inside portholes and welding chamber was taken out by wire electrical discharge machining. As indicated in Fig. 2, the welding zone of extruded profile is formed by the material close to bridge surface, while the matrix zone is formed by the material far away from bridge surface. Thus, the microstructure evolution along two streamlines was focused in this study. The streamline near welding plane is named as welding zone line, where six zones of X1-X6 were determined. X1 located in the lower part of the billet. X2 and X3 corresponded to the upper and lower positions of the vertical side of the bridge. X4 and X5 located in the welding chamber, and X6 was in the position of the die bearing. Thus, the whole porthole die process is included, such as the splitting (X1-X2), welding (X3-X6) and extruding (X6) stages. Meanwhile, the other streamline is named as matrix zone line, and the zones of Z1, Z2 and Z3 were determined. Z1 located inside the billet, and Z2 was in the lower part of the bridge, and Z3 was in the position of the die bearing. The selected zones of X1-X6 and Z1-Z3 were employed for microstructure analysis using electron back scattered diffraction (EBSD) and scanning electron microscopy (SEM). The specimens for EBSD analysis were firstly mechanically polished, and then electro polished by the solution of 20 ml perchloric acid and 80 ml ethanol at the condition of 15 V and 0.6 A for 30 s. In EBSD results, the high angle grain boundaries (HGBs) are defined as the misorientation angle higher than 15°, while the low angle grain boundaries (LGBs) are defined as the misorientation angle within the range of 2–15°.
particles. On the other hand, porthole die is widely used to produce hollow Al profiles with complex cross-section [18]. During porthole die extrusion, the Al billet is firstly divided into several metal streams. Then, these metal streams are contacted to each other and solid bonded inside the welding chamber under the condition of severe pressure and high temperature. Finally, several longitudinal weld seams are formed at the whole length of extruded profile along extrusion direction (ED). Some experimental works have been performed to study the Al profile extruded by porthole die. Bai et al. [19] physically simulated the porthole die extrusion using AA6082 alloy, and the results showed that the degree of local DRX became higher under certain extrusion condition, leading to the reduction of mean grain size. den Bakker et al. [20] found that the mechanical properties of a hollow rectangle Al profile were strongly dependent on the number of the oxide particles located at boundaries of transverse welds. The microstructure and its evolution during the porthole die extrusion should be more complicated, in comparison with flat die extrusion. Firstly, the material flowing route becomes more complex, and the distribution of deformation is more inhomogeneous, which result in the variation of strain, stress and temperature in different zones. Secondly, the longitudinal weld seams are formed due to the occurrence of solid bonding. The material in welding zone experiences the compression, tensile, and shearing stresses, and the deformation degree is quite high. However, the stress condition in matrix zone is relatively simple, and the strain is also at low level. As is known, the occurrence and fraction of DRX are strongly related to deformation condition. Thus, the above facts indicate that Al profiles extruded by porthole die should have significant microstructure inhomogeneity. In our recent study [21], porthole die extrusion was conducted on Al–Zn–Mg alloys, and the extruded Al profile consisted of welding zone and matrix zone. The complete DRX occurred in welding zone, resulting in the formation of fine equiaxed grains. However, large amount of coarse and elongated grains existed in matrix zone due to the occurrence of partial DRX. Although the microstructure inhomogeneity has been found, the dynamic evolution of microstructure during porthole die extrusion is still not clarified. The mechanism of microstructure inhomogeneity should be comprehensively analyzed. Hence, the experiment of porthole die extrusion was performed at temperature of 783 K, using homogenized AA7075 as the billet. After that, the reserved materials inside die cavities such as portholes and welding chamber were taken out. Two streamlines corresponding to welding and matrix zones of the extruded Al profile were selected, and the microstructure of nine zones were carefully examined. The grain structure, texture orientation, and second phase along the streamlines of welding and matrix zones were respectively investigated. Through this study, it is aimed to clarify the microstructure evolution of extruded Al–Zn–Mg profile, which is an important issue for porthole die extrusion.
3. Welding zone characterization 3.1. Grain structure in welding zone Fig. 3 shows the grain structure of materials at X1-X6 along the streamline of welding zone. The grain size of X1 is around 135.8 μm, which is larger than that of the homogenized billet. It indicates that the strain of X1 is relatively low. X2 located at the upper side of the bridge, and it suffered strong shearing deformation during the extrusion process [22,23]. Therefore, the grains of X2 were significantly elongated and some fine grains were formed, which indicated the occurrence of partial DRX with the volume fraction of 30.24%. When the material flowed to X3, complete DRX occurred and fine grain structure was obtained. Then, the material flowed across welding chamber along the path of welding plane (X4-X6), and the stage of solid bonding took place. As is seen, the grain size of X4, X5 and X6 are 1.69 μm, 1.62 μm, 1.71 μm, respectively. Moreover, X4-X6 have the obvious banded grain structure parallel to ED direction due to combined effects of tensile stress along ED direction and compression stress along TD direction. It is noted that the banded grain structure of X3 has deviated angle with ED direction due to combined effects of severe shearing stress provided
2. Experimental procedures The AA7075 ingot with the diameter of ϕ300 mm was casted by semi-continuous method. The chemical compositions of the as-received alloy are listed in Table 1. The homogenization was performed at 803 K for 20 h with air cooling. The orientation map, second phase and pole figures of the homogenized Al–Zn–Mg alloy are shown in Fig. 1. The homogenized billet mainly consisted of equiaxed grains with the average size of 90.0 μm, as shown in Fig. 1(a). Moreover, a large Table 1 Chemical compositions of the as-received Al–Zn–Mg alloy. Element
Zn
Mg
Cu
Cr
Fe
Mn
Ti
Si
Al
wt.%
6.97
2.45
1.31
0.21
0.31
0.05
0.05
0.14
Bal.
2
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Fig. 1. Microstructure characteristics of homogenized Al–Zn–Mg alloy. (a) Grain morphology, (b) second phase distribution, and (c) pole figures.
extruding stages (X3-X6). It is considered that the increasing strain from X1-X2 enhances the occurrence and degree of continuous dynamic recrystallization (CDRX) with gradual transformation from LABs to HABs. This kind of phenomenon was also reported by Kaibyshev et al. [25] and Sakai et al. [26]. At initial stage of CDRX, large amount of dislocations gradually disappear in pairs and surround the substructure merging into a complete crystal structure. The substructure continuously evolves and merges, and fine DRXed grains are formed. After reaching complete CDRX, the fractions of LABs and HABs keep steady [27,28]. Thus, it is known that the fraction of LABs was obviously decreased during splitting stage due to the occurrence of CDRX. However, when the material flowed to X3, complete DRX was achieved, and the fraction of LABs kept steady from X4 to X6. 3.2. Texture evolution in welding zone The inverse pole figures along welding zone line are presented in Fig. 6. The < 111 > concentration along ED with low density of 2.52 can be observed at X1 zone. In case of X2, the concentration of < 101 > along ED and < 111 > along ND were formed due to the constantly shearing deformation. Then, from X2 to X3, the material flowed along bridge surface and rotated around ND for 30°. Consequently, the < 101 > concentration along TD was formed, and the orientation concentration along ED was changed. For the material at zones X4X6, < 111 > and < 001 > orientations along TD were formed due to the tensile stress on welding plane, and the density gradually increased from 4.94 of X4 to 7.36 of X6. The variation of grain orientation indicates the significant evolution of micro-texture during porthole die extrusion. Fig. 7 shows the ODF sections with φ2 of 0°, 45°, 65° along welding zone line. The Miller indices and Euler angles of main texture components are listed in Table 2. There is no significant texture component at X1, and the texture intensity has relatively low value of 1.89. Then, the typical texture component of {111} < 101 > was formed at X2 due to the strong shearing deformation. As mentioned above, when material flowed from X2 to X3, it rotated around ND for around 30°.
Fig. 2. Extrusion setup with the main dimension, and locations of the zones for microstructure observation. TD is the transverse direction, and ND is the normal direction. (Unit: mm).
by bridge and tensile stress along the ED direction. Besides that, the color distribution of zones X1-X6 implies that the grain orientation significantly evolved. The orientation of banded structure with elongated grains is more centralized from X1 to X2, which means the material underwent more concentrated unitary stress at splitting stage. From X3 to X6, < 111 > orientation gradually transformed into < 101 > orientation due to the welding stress provided by welding chamber [24]. Additionally, the grain orientation becomes more uniform from X3 to X6. The average grain size of zones X1-X6 is summarized in Fig. 4, and it confirms that the grains are obviously refined from X1 to X2, while the grain size keeps almost steady from X3 to X6. Fig. 5 presents the misorientation angle of X1-X6 along welding zone line. As is seen, the fraction of LABs (f) significantly decreased from 89.28% (X1) to 61.46% (X2) during splitting stage, and then it was kept at low level and almost steady during subsequent welding and 3
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Fig. 3. Grain Morphology and orientation distribution along welding zone line.
and finally transferred into {101} < 111 > and {101} < 001 > orientations. Fig. 8 schematically illustrates the texture evolution path along the streamline of welding zone. Fig. 9 shows the variation tendency of main texture components based on ODF sections. According to the previous study [29], {112} < 111 > orientation is a typical deformed texture. As is seen from Fig. 9, the proportion of {112} < 111 > drastically increases from 6.7% of X2 to 32.4% of X3, and it keeps at high level from X3 to X6. The increasing of {112} < 111 > indicates that the material suffers severe strain during welding and extruding stages. {101} < 001 > orientation continuously increases from 6.7% of X3 to 17.7% of X6, and {101} < 111 > continuously increases from 16.3% of X3 to 33.3% of X6. As reported by Refs. [30–32], the increasing intensity of {101} < 001 > should be originated from the occurrence of DRX in this study. The proportion of {111} < 101 > reaches the maximum value of 32.2% at X2, and it dramatically decreases to 6.87% of X3 and 1.2% of X6. The proportion of {111} < 112 > increases rapidly from 3.22% of X1 to 38.9% of X4, and then it gradually decreases to 15.5% of X6. The material suffered severe shearing deformation in {111} < 112 > orientation, which is a main component of γ-fiber texture [33]. Moreover, the transformation between {111} < 112 > and {101} < 001 > orientations can also be initiated due to the shearing deformation [34,35]. Up to the final extruding stage, the main textures in welding zone of extruded Al–Zn–Mg profile are {112} < 111 > , {101} < 111 > , {101} < 001 > and {111} < 112 > orientations.
Fig. 4. Average grains size of zones X1-X6 along welding zone line.
Accordingly, the component of {111} < 101 > transferred into {111} < 112 > . Moreover, the texture intensity significantly increased from X2 to X3 because of the accumulated shearing deformation along bridge surface. Since X4-X6 locates on the welding plane, it suffered compression stress along TD, resulting in the grain rotation around TD axis. Consequently, some of the grains rotated along the negative direction of Φ axis and transferred into {112} < 111 > orientation. Another part of the grains rotated along the positive direction of Φ axis 4
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Fig. 5. Relative frequency of misorientation angle of X1 to X6 along welding zone line.
of compression and tensile stress from extrusion die. Fig. 12 shows the grain structure of zones Z1-Z3. As is seen, the grain is significantly elongated along ED direction due to the tensile stress, and the elongating degree was enhanced from Z1 to Z3. It can also be observed that small amount of fine grains existed at the elongated grain boundary. The volume fractions of DRX at Z1, Z2 and Z3 are 1.42%, 1.56% and 1.65%, respectively. The above facts indicate that partial DRX occurred, and DRX degree gradually increased along the welding zone line. Overall, the DRX degree of Z1-Z3 is much lower than that of the material along welding zone line. This should mainly result from the fact that the materials along welding and matrix zones experienced different levels of strain and stress. The material in welding zone experiences serve plastic deformation, while the flowing route of matrix zone is much smoother. Moreover, the width of elongated grain becomes smaller from Z1 to Z3, which indicates the strain gradually increases for the material along matrix zone line. Fig. 11 shows the variation of grain size from Z1 to Z3. As is seen, the grain refinement is a continuous process, and the refining degree is obviously smaller than that of welding zone. Finally, the orientation of elongated grains remains stable from Z1 to Z3 due to the simple strain and stress condition. Fig. 13 shows the misorientation angle at zones Z1-Z3. It can be seen that the fractions of LABs at Z1, Z2 and Z3 are 85.56%, 89.69%, and 84.45%, respectively. As shown in Fig. 11, the coarse elongated grains
3.3. Second phase distribution in welding zone During hot extrusion process, the variation of second phases of Al alloy strongly relates to the complex strain condition and high deformation temperature [36]. Fig. 10 shows the SEM image of second phase distribution of zones X1-X6 along welding zone line. Comparing zone X1 and homogenized billet shown in Fig. 1(b), MgZn2 phase almost disappears, while coarse Al23CuFe4 phase remains unchanged. It should be resulted from the facts that the extrusion temperature of 783 K is sufficient for dissolving MgZn2 phase, but it is insufficient for the dissolution of Al23CuFe4 phase. When the material flowed to X2, the coarse Al23CuFe4 was broken into relatively small pieces, and large amount of MgZn2 precipitated. For zones X3 and X4, the morphology and distribution of Al23CuFe4 is almost as same as those of X2. Moreover, the number of MgZn2 phase at zones X3 and X4 is higher than that at zone X2. 4. Matrix zone characterization 4.1. Grain structure in matrix zone The flowing route of material along matrix zone line is similar with that of flat die extrusion, and the material undergoes the combinations 5
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Table 2 Miller indices and Euler angles for the main texture components. Miller indices
{111} < 101 > {111} < 112 > {112} < 111 > {101} < 111 > {101} < 001 > {001} < 001 >
Euler angles φ1
Φ
φ2
60 90 90 55 0 0
55 55 35 45 45 0
45 45 45 0 0 0
Fig. 8. Diagram of texture evolution along welding zone line during porthole die extrusion.
Fig. 6. Inverse pole figure of X1-X6 along welding zone line.
Fig. 9. The variation of main texture components of X1-X6 along welding zone line.
substructure such as LABs is continuously formed inside the original coarse grains. The original grains are divided by LABs and the proportion of HABs is quite small [37]. Moreover, since the flowing route of matrix zone line is relatively smooth, DRV is the main recovery mechanism for Al–Zn–Mg alloy. Thus, from the viewpoint of misorientation angle, it confirms that the materials in matrix zone have higher DRV degree and lower DRX degree in comparison with welding zone.
4.2. Texture evolution in matrix zone Fig. 7. ODF section of φ2 = 45°, 65° and 90° of X1-X6 along welding zone line.
Fig. 14 shows the inverse pole figures of zones Z1-Z3. The orientation concentration of < 001 > and < 111 > along ED with relatively low density can be found at Z1, and it should be resulted from the tensile stress during extrusion process. When the material flowed to Z2, the continuous tensile stress led to the increasing density of < 001 > and < 111 > orientations along ED, and the orientation concentration
of Z1-Z3 have many substructures, since the material along matrix zone line suffered less stress and strain. Hence, the LABs fraction of matrix zone is obviously much higher than that of welding zone. DRV and slight DRX took place along matrix zone line. During DRV process, the 6
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Fig. 10. Second phase distribution of X1-X4 along welding zone line.
Fig. 11. Orientation distribution figure of Z1-Z3 along matrix zone line.
specifically, from Z2-Z3, {001} < 001 > texture transferred into {112} < 111 > , and {101} < 001 > textures transferred into {101} < 111 > texture. Fig. 16 shows the evolution path of texture component along matrix zone line. Comparing Figs. 8 and 16, it is known that the texture evolution along welding zone line is much more complicated than that along matrix zone line. This difference should be caused by the slight shearing action along the matrix zone line. Fig. 17 shows the proportion variation of main texture components along matrix zone line. As is seen, the proportion of {112} < 111 > decreases from 12.5% of Z1 to 8.61% of Z2, and then it drastically increases to 38.2% of Z3. Similarly, {101} < 111 > decreases from 15.8% of Z1 to 9.5% of Z2, and then it drastically increases to 38.4% of Z3. The proportion vitiation of {001} < 001 > and {112} < 111 > are opposite because of the texture transformation mentioned above. According to the previous studies [38,39], the main
of < 101 > along TD and ND also appeared. Then, when the material flowed to Z3 located at the exit of lower die, < 111 > mainly concentrated along ED, while there is no obvious orientation preference along TD. The orientation variation denotes the texture evolution of matrix zone during extrusion process. Moreover, the density of grain orientation is ever-increasing due to continuously increasing strain from Z1-Z3 along matrix zone line, which is in accordance with above description. Fig. 15 presents the ODF sections of zones Z1-Z3. Z1 exhibits strong {101} < 001 > texture and weak {001} < 001 > texture due to the tensile stress, which agrees well with the results of inverse pole figures. At zone Z2, the type of texture component remains unchanged, while the texture intensity significantly increases. Then, the material suffered complex stress inside welding chamber when it flowed along Z2-Z3, and the type of texture component was finally changed at Z3. More 7
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Fig. 12. Average grains size of Z1-Z3 along matrix zone line. Fig. 14. Inverse pole figure of Z1-Z3 along matrix zone line.
components formed during hot deformation process can be classified as deformed texture and recrystallized texture. {001} < 001 > orientation is a typical recrystallized texture, while {112} < 111 > orientation is a typical deformed texture. With increasing strain along the matrix zone line, {001} < 001 > was gradually transformed into {112} < 111 > [32]. Moreover, the decrease of {001} < 001 > orientation might be because of high DRV degree and low DRX degree in matrix zone. The variation of {101} < 001 > and {101} < 111 > orientations is more related to the tensile stress, which is in accordance with the analysis about inverse pole figure. For {101} < 211 > orientation, its proportion continuously increases from 0.59% of Z1 to 22.2% of Z3. It has been reported that lots of dislocations are generated and slip along the sliding surface due to severe plastic deformation, and the shearing plays an important role in the later stage of the development of {101} < 211 > orientation [34–36]. Up to the final extruding stage, the main textures in matrix zone of extruded Al–Zn–Mg profile are {112} < 111 > , {101} < 111 > and {101} < 211 > orientations.
Fig. 15. ODF section of φ2 = 45°, 65° and 90° at Z1-Z3 along matrix zone line.
Fig. 13. Relative frequency of misorientation angle of Z1-Z3 along matrix zone line. 8
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4.3. Second phase distribution in matrix zone Fig. 18 shows the SEM image of second phase distribution along matrix zone line. For zone Z1 located inside container, the material experienced relatively small strain, and coarse Al23CuFe4 phase still distributes at the triangle area of grain boundaries, while the MgZn2 phase disappears. When the material flowed to Z2, Al23CuFe4 phase was broken into small pieces, and distributed along the elongated grain boundaries. Moreover, a small amount of MgZn2 phases precipitated and distributed randomly. In this study, the extrusion ratio is around 10.5, which is a relatively low value. It is possible that the heat generated by plastic deformation was not sufficient to balance the heat loss, and the temperature of material decreased with the proceeding of extrusion, as reported in Ref. [40]. Consequently, the decrease of temperature caused the precipitation of MgZn2 phase. At zone Z3, the morphology and distribution of Al23CuFe4 phase were not changed, while the number of MgZn2 phase greatly increases. Hence, it further proved that MgZn2 firstly dissolved into Al-matrix, and then gradually precipitated with the proceeding of extrusion. This process should be complete at the lower side of the bridge, and maintained in the final profile. For Al23CuFe4, it cannot dissolve into Al-matrix, while it was gradually broke into small pieces due to the effect of mechanical deformation.
Fig. 16. Diagram of texture evolution along matrix zone line during porthole die extrusion.
5. Conclusion The experiment of porthole die extrusion was conducted using Al–Zn–Mg alloy. The dynamic evolution of microstructure along welding zone line and matrix zone line was analyzed based on EBSD and SEM analysis. The conclusions are drawn as follows. (1) The material of welding zone experienced severe plastic deformation, while the flowing route of matrix zone was relatively smooth. Hence, after the finish of splitting stage, welding zone achieved near complete DRX, and the fraction of LABs keeps at low level. The elongated grains with small amount of fine equiaxed grains were observed along the whole matrix zone line due to the occurrence DRV and slight DRX, and the fraction of LABs keeps at high level. (2) In comparison with the matrix zone, the micro-texture evolution of
Fig. 17. The variation of main texture components of Z1-Z3 along matrix zone line.
Fig. 18. Second phase distribution of Z1-Z3 along matrix zone line. 9
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the welding zone was much more complicated due to the complex compression, shearing and tension stresses. The welding zone of the final extruded profile had the main textures of {112} < 111 > , {101} < 111 > , {101} < 001 > and {111} < 112 > orientations, while the matrix zone of the profile consisted of {112} < 111 > , {101} < 111 > and {101} < 211 > orientations. (3) Along streamlines of both welding and matrix zones, MgZn2 phase firstly dissolved into Al-matrix at the initial extrusion stage, and then it gradually precipitated with an uniform distribution. Al23CuFe4 phase could not dissolve into Al-matrix due to the limitation of extrusion temperature, while it was broke into small pieces due to the effect of mechanical deformation.
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Zhang, Microstructure and texture evolution during porthole die extrusion of Mg-Al-Zn alloy, J. Mater. Process. Technol. 259 (2018) 346–352. [29] K. Hamad, H.W. Yang, Y.G. Ko, Interpretation of annealing texture changes of severely deformed Al-Mg-Si alloy, J. Alloy. Comp. 687 (2016) 300–305. [30] L.F. Shuai, T.L. Huang, G.L. Wu, X. Huang, O.V. Mishin, Development of Goss texture in Al-0.3% Cu annealed after heavy rolling, J. Alloy. Comp. 749 (2018) 399–405. [31] M.R. Toroghinejad, F. Ashrafizadeh, R. Jamaati, M. Hoseini, J.A. Szpunar, Textural evolution of nanostructured AA5083 produced by ARB, Mater. Sci. Eng. A 556 (2012) 351–357. [32] M.R. Toroghinejad, R. Jamaati, M. Hoseini, J.A. Szpunar, J. Dutkiewicz, Texture Evolution of nanostructured aluminum/copper composite produced by the accumulative roll bonding and folding process, Metall. Mater. Trans. A 44 (2013) 1587–1598. [33] R. Jamaati, M.R. Toroghinejad, Effect of stacking fault energy on deformation texture development of nanostructured materials produced by the ARB process, Mater. Sci. Eng. A 598 (2014) 263–276. [34] F. Montheillet, M. Cohen, J.J. Jonas, Axial stresses and texture development during the torsion testing of Al, Cu and Alpha-Fe, Acta Metall. 32 (1984) 2077–2089. [35] S. Li, I.J. Beyerlein, M.A.M. Bourke, Texture formation during equal channel angular extrusion of fcc and bcc materials: comparison with simple shear, Mater. Sci. Eng. A 394 (2005) 66–77. [36] G. Chen, L. Chen, G. Zhao, C.S. Zhang, Microstructure evolution during solution treatment of extruded Al-Zn-Mg profile containing a longitudinal weld seam, J. Alloy. Comp. 729 (2017) 210–221. [37] H. Asgharzadeh, A. Simchi, H.S. Kim, Dynamic restoration and microstructural evolution during hot deformation of a P/M Al6063 alloy, Mater. Sci. Eng. A 542 (2012) 56–63. [38] T. Guo, J.S. Li, J. Wang, W.Y. Wang, Y. Liu, X.M. Luo, H.C. Kou, E. 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Data availability The raw/processed data required to reproduce these findings cannot be shared at this time due to technical or time limitations. Acknowledgements The authors would like to acknowledge the financial support from National Natural Science Foundation of China (U1708251, 51735008), and Key Research and Development Program of Shandong Province (2018GGX103041), and Young Scholars Program of Shandong University (2018WLJH26). References [1] Q. Ding, D. Zhang, J. Zuo, S. Hou, L. Zhuang, J. Zhang, The effect of grain boundary character evolution on the intergranular corrosion behavior of advanced Al-Mg-3wt % Zn alloy with Mg variation, Mater. Char. 146 (2018) 47–54. [2] T. Aoba, M. Kobayashi, H. Miura, Effects of aging on mechanical properties and microstructure of multi-directionally forged 7075 aluminum alloy, Mater. Sci. Eng. A 700 (2017) 220–225. [3] X. Wang, Q. Pan, L. Liu, S. Xiong, W. Wang, J. Lai, Y. Sun, Z. Huang, Characterization of hot extrusion and heat treatment on mechanical properties in a spray formed ultra-high strength Al-Zn-Mg-Cu alloy, Mater. Char. 144 (2018) 131–140. [4] L. Lin, Z. Liu, P. Ying, M. Liu, Improved stress corrosion cracking resistance and strength of a two-step aged Al-Zn-Mg-Cu alloy using taguchi method, J. Mater. Eng. Perform. 24 (2015) 4870–4877. [5] I. Balasundar, M.S. Rao, T. Raghu, Equal channel angular pressing die to extrude a variety of materials, Mater. Des. 30 (2009) 1050–1059. [6] M.A. Afifi, Y.C. Wang, X. Cheng, S. Li, T.G. Langdon, Strain rate dependence of compressive behavior in an Al-Zn-Mg alloy processed by ECAP, J. Alloy. Comp. 791 (2019) 1079–1087. [7] H. Ghandvar, M.H. Idris, N. Ahmad, Effect of hot extrusion on microstructural evolution and tensile properties of Al-15%Mg2Si-xGd in-situ composites, J. Alloy. Comp. 751 (2018) 370–390. [8] F. Gagliardi, C. Ciancio, G. Ambrogio, Optimization of porthole die extrusion by grey-taguchi relational analysis, Int. J. Adv. Manuf. Technol. 94 (2018) 719–728. [9] X. Fan, L. Chen, G. Chen, G. Zhao, C. Zhang, Joining of 1060/6063 aluminum alloys based on porthole die extrusion process, J. Mater. Process. Technol. 250 (2017) 65–72. [10] A. Mogucheva, E. Babich, B. Ovsyannikov, R. Kaibyshev, Microstructural evolution in a 5024 aluminum alloy processed by ECAP with and without back pressure, Mater. Sci. Eng. A 560 (2013) 178–192. [11] M. Moghaddam, A. Zarei-Hanzaki, M.H. Pishbin, A.H. Shafieizad, V.B. Oliveira, Characterization of the microstructure, texture and mechanical properties of 7075 aluminum alloy in early stage of severe plastic deformation, Mater. Char. 119 (2016) 137–147. [12] P. Shaterani, A. Zarei-Hanzaki, S.M. Fatemi-Varzaneh, S.B. Hassas-Irani, The second phase particles and mechanical properties of 2124 aluminum alloy processed by accumulative back extrusion, Mater. Des. 58 (2014) 535–542. [13] X. Dong, F. Chen, S. Chen, Y. Liu, Z. Huang, H. Chen, S. Feng, L. Zhao, Z. Wu, X. Zhang, Microstructure and microhardness of hot extruded 7075 aluminum alloy micro-gear, J. Mater. Process. Technol. 219 (2015) 199–208. [14] A. Bois-Brochu, C. Blais, F.A.T. Goma, D. Larouche, J. Boselli, M. Brochu, Characterization of Al-Li 2099 extrusions and the influence of fiber texture on the
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Evolution of grain structure, micro-texture and second phase during porthole die extrusion of Al–Zn–Mg alloy
T
Liang Chen, Gaojin Chen, Jianwei Tang, Guoqun Zhao∗, Cunsheng Zhang Key Laboratory for Liquid-Solid Structural Evolution and Processing of Materials (Ministry of Education), Shandong University, Jinan, Shandong, 250061, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: Al–Zn–Mg alloy Extrusion Grain structure Texture Second phase
The porthole die extrusion of high strength Al–Zn–Mg alloy was carried out at 783 K. The reserved materials inside portholes and welding chamber were taken out to investigate the dynamic evolution of grain structure, micro-texture and second phase along streamlines of welding zone and matrix zone. The material of welding zone experienced severe plastic deformation, while the flowing route of matrix zone was relatively smooth. Hence, the welding zone achieved near complete dynamic crystallization after splitting stage, and the grain size kept steady during the subsequent welding and extruding stage. The streamline of matrix zone mainly consisted of elongated grains with small amount of fine equiaxed grains, which indicated the occurrence of dynamic recovery and slight dynamic crystallization. The texture evolution of welding zone was much more complicated due to the combination of compression, shearing and tension stresses. The welding zone of the final extruded profile had the main textures of {112} < 111 > , {101} < 111 > , {101} < 001 > and {111} < 112 > orientations, while the matrix zone of the profile consisted of {112} < 111 > , {101} < 111 > and {101} < 211 > orientations. The extrusion temperature is sufficient for dissolving MgZn2 phase, while it is insufficient for the dissolution of Al23CuFe4. Instead, the coarse Al23CuFe4 was broke into small pieces due to the effects of mechanical deformation.
1. Introduction High strength Al–Zn–Mg alloys have been widely applied in the fields of aerospace, high speed train and automobile, due to their advantages of low density and high specific strength [1–4]. Hot extrusion is one of the main forming processes on Al alloys to produce extruded profiles with identical cross-section. In order to obtain Al profiles with high dimensional accuracy and excellent mechanical properties, both the material flowing behavior and microstructure evolution during hot extrusion should be well controlled. In the past decade, the researchers have attempted to enhance the flowing homogeneity through optimizing the die structure and process parameters [5–8]. However, the evolution of microstructure during hot extrusion has not been clarified. As is known, the grain morphology, texture feature and second phase evolve markedly during hot extrusion process. The occurrence of dynamic recovery (DRV) and dynamic recrystallization (DRX) can help to reduce the grain size, and it is beneficial for improving the mechanical properties of extruded profiles [9,10]. Moreover, the texture evolution such as the type, density, and distribution also plays an important role on the final mechanical properties. It has been proved that the anisotropy of extruded Al profiles can be obviously enhanced by
∗
controlling the texture during hot extrusion [11]. In case of second phases, their composition, morphology and distribution should be well examined. In general, fine second phase with uniform and dispersed distribution can effectively improve the microstructure, while coarse second phase has adverse effects [12,13]. Flat die is usually used in practical production to produce solid Al profiles with simple cross-section. The researchers have carried out lots of experimental works to study the microstructure characteristics of solid Al profiles. Bois-Brochu et al. [14] examined the microstructure, texture and static mechanical properties of extruded Al–Li alloy, and found that the variation of strength and anisotropy was strongly dependent on billet temperature. Xu et al. [15] proposed that the size of eutectic Si particles and primary intermetallic particles of Al–Si–Cu alloy was reduced after hot extrusion, and the profile extruded at 450 °C exhibited the best combination of strength and ductility. Li et al. [16] proposed that the asymmetric feeder chamber can modify the grain structure, texture component and second phase distribution of the extruded Al–Zn–Mg alloy, which led to the decreasing anisotropy of elongation. Wu et al. [17] conducted extrusion experiment using Al–Si–Mg alloys with various Si contents, and found that the DRX behavior during hot extrusion was strongly related to the number of large Si
Corresponding author. E-mail address: [email protected] (G. Zhao).
https://doi.org/10.1016/j.matchar.2019.109953 Received 27 June 2019; Received in revised form 14 August 2019; Accepted 1 October 2019 Available online 16 October 2019 1044-5803/ © 2019 Elsevier Inc. All rights reserved.
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amount of second phases can be observed from Fig. 1(b). EDS data implies that the second phases located at triangle region of grain boundary are Al23CuFe4, while those distributed inside the grains are MgZn2. In addition, there are also some second phases, which are characterized as Al2CuMg, discontinuously distribute along grain boundary. As shown in Fig. 1(c), the peak texture intensity is as low as 2.77, which means that the homogenized alloy has no obvious orientation preference. The extrusion setup and the main dimension of the bridge, upper die, and lower die are schematically drawn in Fig. 2. The bridge is separated from upper die in order to facilitate the removal of reserved material inside portholes and welding chamber after extrusion experiment. Moreover, for convenience of taking specimens for microstructure observation, a plate shaped profile was extruded out. The homogenized billet with the dimension of ϕ40 × 60 mm was machined. Then, the billet was firstly heated to 783 K together with the die, and held for 20 min. After that, the extrusion was started, and the ram velocity is set as 6 mm/min. The experiment was stopped when the stroke reached 30 mm, and the whole setup was quenched by cooling water. Finally, the remaining materials inside portholes and welding chamber was taken out by wire electrical discharge machining. As indicated in Fig. 2, the welding zone of extruded profile is formed by the material close to bridge surface, while the matrix zone is formed by the material far away from bridge surface. Thus, the microstructure evolution along two streamlines was focused in this study. The streamline near welding plane is named as welding zone line, where six zones of X1-X6 were determined. X1 located in the lower part of the billet. X2 and X3 corresponded to the upper and lower positions of the vertical side of the bridge. X4 and X5 located in the welding chamber, and X6 was in the position of the die bearing. Thus, the whole porthole die process is included, such as the splitting (X1-X2), welding (X3-X6) and extruding (X6) stages. Meanwhile, the other streamline is named as matrix zone line, and the zones of Z1, Z2 and Z3 were determined. Z1 located inside the billet, and Z2 was in the lower part of the bridge, and Z3 was in the position of the die bearing. The selected zones of X1-X6 and Z1-Z3 were employed for microstructure analysis using electron back scattered diffraction (EBSD) and scanning electron microscopy (SEM). The specimens for EBSD analysis were firstly mechanically polished, and then electro polished by the solution of 20 ml perchloric acid and 80 ml ethanol at the condition of 15 V and 0.6 A for 30 s. In EBSD results, the high angle grain boundaries (HGBs) are defined as the misorientation angle higher than 15°, while the low angle grain boundaries (LGBs) are defined as the misorientation angle within the range of 2–15°.
particles. On the other hand, porthole die is widely used to produce hollow Al profiles with complex cross-section [18]. During porthole die extrusion, the Al billet is firstly divided into several metal streams. Then, these metal streams are contacted to each other and solid bonded inside the welding chamber under the condition of severe pressure and high temperature. Finally, several longitudinal weld seams are formed at the whole length of extruded profile along extrusion direction (ED). Some experimental works have been performed to study the Al profile extruded by porthole die. Bai et al. [19] physically simulated the porthole die extrusion using AA6082 alloy, and the results showed that the degree of local DRX became higher under certain extrusion condition, leading to the reduction of mean grain size. den Bakker et al. [20] found that the mechanical properties of a hollow rectangle Al profile were strongly dependent on the number of the oxide particles located at boundaries of transverse welds. The microstructure and its evolution during the porthole die extrusion should be more complicated, in comparison with flat die extrusion. Firstly, the material flowing route becomes more complex, and the distribution of deformation is more inhomogeneous, which result in the variation of strain, stress and temperature in different zones. Secondly, the longitudinal weld seams are formed due to the occurrence of solid bonding. The material in welding zone experiences the compression, tensile, and shearing stresses, and the deformation degree is quite high. However, the stress condition in matrix zone is relatively simple, and the strain is also at low level. As is known, the occurrence and fraction of DRX are strongly related to deformation condition. Thus, the above facts indicate that Al profiles extruded by porthole die should have significant microstructure inhomogeneity. In our recent study [21], porthole die extrusion was conducted on Al–Zn–Mg alloys, and the extruded Al profile consisted of welding zone and matrix zone. The complete DRX occurred in welding zone, resulting in the formation of fine equiaxed grains. However, large amount of coarse and elongated grains existed in matrix zone due to the occurrence of partial DRX. Although the microstructure inhomogeneity has been found, the dynamic evolution of microstructure during porthole die extrusion is still not clarified. The mechanism of microstructure inhomogeneity should be comprehensively analyzed. Hence, the experiment of porthole die extrusion was performed at temperature of 783 K, using homogenized AA7075 as the billet. After that, the reserved materials inside die cavities such as portholes and welding chamber were taken out. Two streamlines corresponding to welding and matrix zones of the extruded Al profile were selected, and the microstructure of nine zones were carefully examined. The grain structure, texture orientation, and second phase along the streamlines of welding and matrix zones were respectively investigated. Through this study, it is aimed to clarify the microstructure evolution of extruded Al–Zn–Mg profile, which is an important issue for porthole die extrusion.
3. Welding zone characterization 3.1. Grain structure in welding zone Fig. 3 shows the grain structure of materials at X1-X6 along the streamline of welding zone. The grain size of X1 is around 135.8 μm, which is larger than that of the homogenized billet. It indicates that the strain of X1 is relatively low. X2 located at the upper side of the bridge, and it suffered strong shearing deformation during the extrusion process [22,23]. Therefore, the grains of X2 were significantly elongated and some fine grains were formed, which indicated the occurrence of partial DRX with the volume fraction of 30.24%. When the material flowed to X3, complete DRX occurred and fine grain structure was obtained. Then, the material flowed across welding chamber along the path of welding plane (X4-X6), and the stage of solid bonding took place. As is seen, the grain size of X4, X5 and X6 are 1.69 μm, 1.62 μm, 1.71 μm, respectively. Moreover, X4-X6 have the obvious banded grain structure parallel to ED direction due to combined effects of tensile stress along ED direction and compression stress along TD direction. It is noted that the banded grain structure of X3 has deviated angle with ED direction due to combined effects of severe shearing stress provided
2. Experimental procedures The AA7075 ingot with the diameter of ϕ300 mm was casted by semi-continuous method. The chemical compositions of the as-received alloy are listed in Table 1. The homogenization was performed at 803 K for 20 h with air cooling. The orientation map, second phase and pole figures of the homogenized Al–Zn–Mg alloy are shown in Fig. 1. The homogenized billet mainly consisted of equiaxed grains with the average size of 90.0 μm, as shown in Fig. 1(a). Moreover, a large Table 1 Chemical compositions of the as-received Al–Zn–Mg alloy. Element
Zn
Mg
Cu
Cr
Fe
Mn
Ti
Si
Al
wt.%
6.97
2.45
1.31
0.21
0.31
0.05
0.05
0.14
Bal.
2
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Fig. 1. Microstructure characteristics of homogenized Al–Zn–Mg alloy. (a) Grain morphology, (b) second phase distribution, and (c) pole figures.
extruding stages (X3-X6). It is considered that the increasing strain from X1-X2 enhances the occurrence and degree of continuous dynamic recrystallization (CDRX) with gradual transformation from LABs to HABs. This kind of phenomenon was also reported by Kaibyshev et al. [25] and Sakai et al. [26]. At initial stage of CDRX, large amount of dislocations gradually disappear in pairs and surround the substructure merging into a complete crystal structure. The substructure continuously evolves and merges, and fine DRXed grains are formed. After reaching complete CDRX, the fractions of LABs and HABs keep steady [27,28]. Thus, it is known that the fraction of LABs was obviously decreased during splitting stage due to the occurrence of CDRX. However, when the material flowed to X3, complete DRX was achieved, and the fraction of LABs kept steady from X4 to X6. 3.2. Texture evolution in welding zone The inverse pole figures along welding zone line are presented in Fig. 6. The < 111 > concentration along ED with low density of 2.52 can be observed at X1 zone. In case of X2, the concentration of < 101 > along ED and < 111 > along ND were formed due to the constantly shearing deformation. Then, from X2 to X3, the material flowed along bridge surface and rotated around ND for 30°. Consequently, the < 101 > concentration along TD was formed, and the orientation concentration along ED was changed. For the material at zones X4X6, < 111 > and < 001 > orientations along TD were formed due to the tensile stress on welding plane, and the density gradually increased from 4.94 of X4 to 7.36 of X6. The variation of grain orientation indicates the significant evolution of micro-texture during porthole die extrusion. Fig. 7 shows the ODF sections with φ2 of 0°, 45°, 65° along welding zone line. The Miller indices and Euler angles of main texture components are listed in Table 2. There is no significant texture component at X1, and the texture intensity has relatively low value of 1.89. Then, the typical texture component of {111} < 101 > was formed at X2 due to the strong shearing deformation. As mentioned above, when material flowed from X2 to X3, it rotated around ND for around 30°.
Fig. 2. Extrusion setup with the main dimension, and locations of the zones for microstructure observation. TD is the transverse direction, and ND is the normal direction. (Unit: mm).
by bridge and tensile stress along the ED direction. Besides that, the color distribution of zones X1-X6 implies that the grain orientation significantly evolved. The orientation of banded structure with elongated grains is more centralized from X1 to X2, which means the material underwent more concentrated unitary stress at splitting stage. From X3 to X6, < 111 > orientation gradually transformed into < 101 > orientation due to the welding stress provided by welding chamber [24]. Additionally, the grain orientation becomes more uniform from X3 to X6. The average grain size of zones X1-X6 is summarized in Fig. 4, and it confirms that the grains are obviously refined from X1 to X2, while the grain size keeps almost steady from X3 to X6. Fig. 5 presents the misorientation angle of X1-X6 along welding zone line. As is seen, the fraction of LABs (f) significantly decreased from 89.28% (X1) to 61.46% (X2) during splitting stage, and then it was kept at low level and almost steady during subsequent welding and 3
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Fig. 3. Grain Morphology and orientation distribution along welding zone line.
and finally transferred into {101} < 111 > and {101} < 001 > orientations. Fig. 8 schematically illustrates the texture evolution path along the streamline of welding zone. Fig. 9 shows the variation tendency of main texture components based on ODF sections. According to the previous study [29], {112} < 111 > orientation is a typical deformed texture. As is seen from Fig. 9, the proportion of {112} < 111 > drastically increases from 6.7% of X2 to 32.4% of X3, and it keeps at high level from X3 to X6. The increasing of {112} < 111 > indicates that the material suffers severe strain during welding and extruding stages. {101} < 001 > orientation continuously increases from 6.7% of X3 to 17.7% of X6, and {101} < 111 > continuously increases from 16.3% of X3 to 33.3% of X6. As reported by Refs. [30–32], the increasing intensity of {101} < 001 > should be originated from the occurrence of DRX in this study. The proportion of {111} < 101 > reaches the maximum value of 32.2% at X2, and it dramatically decreases to 6.87% of X3 and 1.2% of X6. The proportion of {111} < 112 > increases rapidly from 3.22% of X1 to 38.9% of X4, and then it gradually decreases to 15.5% of X6. The material suffered severe shearing deformation in {111} < 112 > orientation, which is a main component of γ-fiber texture [33]. Moreover, the transformation between {111} < 112 > and {101} < 001 > orientations can also be initiated due to the shearing deformation [34,35]. Up to the final extruding stage, the main textures in welding zone of extruded Al–Zn–Mg profile are {112} < 111 > , {101} < 111 > , {101} < 001 > and {111} < 112 > orientations.
Fig. 4. Average grains size of zones X1-X6 along welding zone line.
Accordingly, the component of {111} < 101 > transferred into {111} < 112 > . Moreover, the texture intensity significantly increased from X2 to X3 because of the accumulated shearing deformation along bridge surface. Since X4-X6 locates on the welding plane, it suffered compression stress along TD, resulting in the grain rotation around TD axis. Consequently, some of the grains rotated along the negative direction of Φ axis and transferred into {112} < 111 > orientation. Another part of the grains rotated along the positive direction of Φ axis 4
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Fig. 5. Relative frequency of misorientation angle of X1 to X6 along welding zone line.
of compression and tensile stress from extrusion die. Fig. 12 shows the grain structure of zones Z1-Z3. As is seen, the grain is significantly elongated along ED direction due to the tensile stress, and the elongating degree was enhanced from Z1 to Z3. It can also be observed that small amount of fine grains existed at the elongated grain boundary. The volume fractions of DRX at Z1, Z2 and Z3 are 1.42%, 1.56% and 1.65%, respectively. The above facts indicate that partial DRX occurred, and DRX degree gradually increased along the welding zone line. Overall, the DRX degree of Z1-Z3 is much lower than that of the material along welding zone line. This should mainly result from the fact that the materials along welding and matrix zones experienced different levels of strain and stress. The material in welding zone experiences serve plastic deformation, while the flowing route of matrix zone is much smoother. Moreover, the width of elongated grain becomes smaller from Z1 to Z3, which indicates the strain gradually increases for the material along matrix zone line. Fig. 11 shows the variation of grain size from Z1 to Z3. As is seen, the grain refinement is a continuous process, and the refining degree is obviously smaller than that of welding zone. Finally, the orientation of elongated grains remains stable from Z1 to Z3 due to the simple strain and stress condition. Fig. 13 shows the misorientation angle at zones Z1-Z3. It can be seen that the fractions of LABs at Z1, Z2 and Z3 are 85.56%, 89.69%, and 84.45%, respectively. As shown in Fig. 11, the coarse elongated grains
3.3. Second phase distribution in welding zone During hot extrusion process, the variation of second phases of Al alloy strongly relates to the complex strain condition and high deformation temperature [36]. Fig. 10 shows the SEM image of second phase distribution of zones X1-X6 along welding zone line. Comparing zone X1 and homogenized billet shown in Fig. 1(b), MgZn2 phase almost disappears, while coarse Al23CuFe4 phase remains unchanged. It should be resulted from the facts that the extrusion temperature of 783 K is sufficient for dissolving MgZn2 phase, but it is insufficient for the dissolution of Al23CuFe4 phase. When the material flowed to X2, the coarse Al23CuFe4 was broken into relatively small pieces, and large amount of MgZn2 precipitated. For zones X3 and X4, the morphology and distribution of Al23CuFe4 is almost as same as those of X2. Moreover, the number of MgZn2 phase at zones X3 and X4 is higher than that at zone X2. 4. Matrix zone characterization 4.1. Grain structure in matrix zone The flowing route of material along matrix zone line is similar with that of flat die extrusion, and the material undergoes the combinations 5
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Table 2 Miller indices and Euler angles for the main texture components. Miller indices
{111} < 101 > {111} < 112 > {112} < 111 > {101} < 111 > {101} < 001 > {001} < 001 >
Euler angles φ1
Φ
φ2
60 90 90 55 0 0
55 55 35 45 45 0
45 45 45 0 0 0
Fig. 8. Diagram of texture evolution along welding zone line during porthole die extrusion.
Fig. 6. Inverse pole figure of X1-X6 along welding zone line.
Fig. 9. The variation of main texture components of X1-X6 along welding zone line.
substructure such as LABs is continuously formed inside the original coarse grains. The original grains are divided by LABs and the proportion of HABs is quite small [37]. Moreover, since the flowing route of matrix zone line is relatively smooth, DRV is the main recovery mechanism for Al–Zn–Mg alloy. Thus, from the viewpoint of misorientation angle, it confirms that the materials in matrix zone have higher DRV degree and lower DRX degree in comparison with welding zone.
4.2. Texture evolution in matrix zone Fig. 7. ODF section of φ2 = 45°, 65° and 90° of X1-X6 along welding zone line.
Fig. 14 shows the inverse pole figures of zones Z1-Z3. The orientation concentration of < 001 > and < 111 > along ED with relatively low density can be found at Z1, and it should be resulted from the tensile stress during extrusion process. When the material flowed to Z2, the continuous tensile stress led to the increasing density of < 001 > and < 111 > orientations along ED, and the orientation concentration
of Z1-Z3 have many substructures, since the material along matrix zone line suffered less stress and strain. Hence, the LABs fraction of matrix zone is obviously much higher than that of welding zone. DRV and slight DRX took place along matrix zone line. During DRV process, the 6
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Fig. 10. Second phase distribution of X1-X4 along welding zone line.
Fig. 11. Orientation distribution figure of Z1-Z3 along matrix zone line.
specifically, from Z2-Z3, {001} < 001 > texture transferred into {112} < 111 > , and {101} < 001 > textures transferred into {101} < 111 > texture. Fig. 16 shows the evolution path of texture component along matrix zone line. Comparing Figs. 8 and 16, it is known that the texture evolution along welding zone line is much more complicated than that along matrix zone line. This difference should be caused by the slight shearing action along the matrix zone line. Fig. 17 shows the proportion variation of main texture components along matrix zone line. As is seen, the proportion of {112} < 111 > decreases from 12.5% of Z1 to 8.61% of Z2, and then it drastically increases to 38.2% of Z3. Similarly, {101} < 111 > decreases from 15.8% of Z1 to 9.5% of Z2, and then it drastically increases to 38.4% of Z3. The proportion vitiation of {001} < 001 > and {112} < 111 > are opposite because of the texture transformation mentioned above. According to the previous studies [38,39], the main
of < 101 > along TD and ND also appeared. Then, when the material flowed to Z3 located at the exit of lower die, < 111 > mainly concentrated along ED, while there is no obvious orientation preference along TD. The orientation variation denotes the texture evolution of matrix zone during extrusion process. Moreover, the density of grain orientation is ever-increasing due to continuously increasing strain from Z1-Z3 along matrix zone line, which is in accordance with above description. Fig. 15 presents the ODF sections of zones Z1-Z3. Z1 exhibits strong {101} < 001 > texture and weak {001} < 001 > texture due to the tensile stress, which agrees well with the results of inverse pole figures. At zone Z2, the type of texture component remains unchanged, while the texture intensity significantly increases. Then, the material suffered complex stress inside welding chamber when it flowed along Z2-Z3, and the type of texture component was finally changed at Z3. More 7
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Fig. 12. Average grains size of Z1-Z3 along matrix zone line. Fig. 14. Inverse pole figure of Z1-Z3 along matrix zone line.
components formed during hot deformation process can be classified as deformed texture and recrystallized texture. {001} < 001 > orientation is a typical recrystallized texture, while {112} < 111 > orientation is a typical deformed texture. With increasing strain along the matrix zone line, {001} < 001 > was gradually transformed into {112} < 111 > [32]. Moreover, the decrease of {001} < 001 > orientation might be because of high DRV degree and low DRX degree in matrix zone. The variation of {101} < 001 > and {101} < 111 > orientations is more related to the tensile stress, which is in accordance with the analysis about inverse pole figure. For {101} < 211 > orientation, its proportion continuously increases from 0.59% of Z1 to 22.2% of Z3. It has been reported that lots of dislocations are generated and slip along the sliding surface due to severe plastic deformation, and the shearing plays an important role in the later stage of the development of {101} < 211 > orientation [34–36]. Up to the final extruding stage, the main textures in matrix zone of extruded Al–Zn–Mg profile are {112} < 111 > , {101} < 111 > and {101} < 211 > orientations.
Fig. 15. ODF section of φ2 = 45°, 65° and 90° at Z1-Z3 along matrix zone line.
Fig. 13. Relative frequency of misorientation angle of Z1-Z3 along matrix zone line. 8
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4.3. Second phase distribution in matrix zone Fig. 18 shows the SEM image of second phase distribution along matrix zone line. For zone Z1 located inside container, the material experienced relatively small strain, and coarse Al23CuFe4 phase still distributes at the triangle area of grain boundaries, while the MgZn2 phase disappears. When the material flowed to Z2, Al23CuFe4 phase was broken into small pieces, and distributed along the elongated grain boundaries. Moreover, a small amount of MgZn2 phases precipitated and distributed randomly. In this study, the extrusion ratio is around 10.5, which is a relatively low value. It is possible that the heat generated by plastic deformation was not sufficient to balance the heat loss, and the temperature of material decreased with the proceeding of extrusion, as reported in Ref. [40]. Consequently, the decrease of temperature caused the precipitation of MgZn2 phase. At zone Z3, the morphology and distribution of Al23CuFe4 phase were not changed, while the number of MgZn2 phase greatly increases. Hence, it further proved that MgZn2 firstly dissolved into Al-matrix, and then gradually precipitated with the proceeding of extrusion. This process should be complete at the lower side of the bridge, and maintained in the final profile. For Al23CuFe4, it cannot dissolve into Al-matrix, while it was gradually broke into small pieces due to the effect of mechanical deformation.
Fig. 16. Diagram of texture evolution along matrix zone line during porthole die extrusion.
5. Conclusion The experiment of porthole die extrusion was conducted using Al–Zn–Mg alloy. The dynamic evolution of microstructure along welding zone line and matrix zone line was analyzed based on EBSD and SEM analysis. The conclusions are drawn as follows. (1) The material of welding zone experienced severe plastic deformation, while the flowing route of matrix zone was relatively smooth. Hence, after the finish of splitting stage, welding zone achieved near complete DRX, and the fraction of LABs keeps at low level. The elongated grains with small amount of fine equiaxed grains were observed along the whole matrix zone line due to the occurrence DRV and slight DRX, and the fraction of LABs keeps at high level. (2) In comparison with the matrix zone, the micro-texture evolution of
Fig. 17. The variation of main texture components of Z1-Z3 along matrix zone line.
Fig. 18. Second phase distribution of Z1-Z3 along matrix zone line. 9
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the welding zone was much more complicated due to the complex compression, shearing and tension stresses. The welding zone of the final extruded profile had the main textures of {112} < 111 > , {101} < 111 > , {101} < 001 > and {111} < 112 > orientations, while the matrix zone of the profile consisted of {112} < 111 > , {101} < 111 > and {101} < 211 > orientations. (3) Along streamlines of both welding and matrix zones, MgZn2 phase firstly dissolved into Al-matrix at the initial extrusion stage, and then it gradually precipitated with an uniform distribution. Al23CuFe4 phase could not dissolve into Al-matrix due to the limitation of extrusion temperature, while it was broke into small pieces due to the effect of mechanical deformation.
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