Fig. 5 Inverse pole figures (a, d, g, j), grain boundary angle distribution figures (b, e, h, k) and twin distribution figures (c, f, i, l) of as-compressed alloy at 200 °C with different strains

In order to further investigate the twin types and DRX mechanisms, the TEM observation of the as-compressed alloy at 200 °C was carried out. Figure 6(b) indicates that double twin lamellae of ~0.7 μm is observed at the strain of 0.15, the type of which is - according to EBSD analysis (Fig. 5(f)) [7,24]. Compared with the relatively high density of dislocations, the higher density of dislocation clusters and rearrangement constitute the dislocation walls in contraction twins at the strain of 0.45, as shown in Fig. 6(c). As mentioned above, the formation of twin bands (a set of parallel and symmetrical  compression twin lamellae) results in ineffective release of stress via formation of contraction twins and double twins. The twin bands continue to trapping dislocations, which can generate a great deal of dislocation walls, second twins, double twins and even micro-twins in the second twins (as shown in Figs. 6(c) and (d)). Moreover, as the strain is increased up to 0.6, a significant number of sub-grains are formed, which is ascribed to the accumulation of stress. The accumulation of stress induces the activation of the non-basal slip ＜a+c＞ dislocations inside the twin bands and around the twin boundaries.

Fig. 6 TEM images of as-compressed alloy at 200 °C with different strains

Fig. 7 Schematic representation of twin DRX mechanism at 200 °C (red parallelogram regions show nuclei)

The above phenomena of these three kinds of twin nucleation processes are recognized as the features of the TDRX mechanism [6,7]. The schematic illustration of the twin DRX nucleation in twin bands is shown in   Fig. 7. As seen in Fig. 7(a), in the twin bands, the dislocations pile up in interiors of twins. The dislocation walls are formed by rearrangement at a certain angle with primary twins, and then the subsequent dislocations assist the walls to migrate locally in order to reduce their stored energy. Consequently, the deformation twin is subdivided into nuclei by the development of transverse low-angled boundaries [6]. Figure 7(b) shows - double twin which is the second  twin occurring in the primary  twin. As reported, the basal slip dislocations move across the  twin and pile up at the  twin boundaries in double twins, resulting in the formation of low-angle boundaries inside or near double twins [7]. The second twins and micro-twins within primary twin lamellae are shown in Fig. 7(c). Similar to Fig. 7(b), the formation of the nuclei is related with both the basal slip dislocations and the non-basal slip ＜a+c＞ dislocations, which are pinned by the twin-walls during the formation of the micro-twin boundaries and the second twin boundaries [25]. The critical resolved shear stress (CRSS) for the activation of the non-basal slip ＜a+c＞ dislocations is so high at room temperature that the distortion is needed to be further accumulated via dislocation pile-ups at a strain of 0.6. The sub-grains can form via three kinds of TDRX mechanisms and gradually transform into random boundaries by absorbing more dislocations due to an additional increase in misorientations within matrix. However, the TDRX grains cannot grow up because of the poor grain boundary migration under the condition of low temperature. Since the grain boundary migration depends on temperature, it is a rate-controlling process. The DRX occurs at 200 °C and exhibits obvious strain rate sensitivity [25]. The shear bands begin to form while the twin bands are consumed via the formation of ultrafine TDRX grains in the present investigation [22]. In other words, the shear bands can be defined as streamlined strips, which are composed of broken twins (compression and double twin lamellae) and tiny TDRX grains.

Figure 8 shows the EBSD maps of the specimens after compression at 350 °C. The microstructure mainly includes twins at a small strain of 0.08, and the type of the twins is  extension twin, as shown in Fig. 8(c). According to the previous research [20], the number of ＜a+c＞non-basal slip dislocation is large at low temperature and drops rapidly as the temperature increases whereas the twinning is relatively insensitive to the temperature. Therefore, compared with that at 200 °C, the number of twins decreases because the plasticity capability is developed by activation of non-basal slips at 350 °C (as seen in Fig. 3(a) and Fig. 4(a)). In addition, some low-angle grain boundaries appear (as seen in Fig. 8(b)), resulting from the accumulation of a large number of ＜a+c＞slip and slip dislocations near the twin boundaries and grain boundaries.

With the development of the deformation degree, the migration ability of twin boundary is improved gradually as well as the interaction between twin boundaries and dislocations, which results in a transformation from twin boundaries into random grain boundaries [26]. Hence, when the strain is up to 0.15, the DRX occurs at both the twin boundaries and initial grain boundaries (Fig. 8(d)). In the meantime, the number of extension twins dwindles by the recrystallization, as shown in Fig. 8(f). It is well indicated that the severer deformation is, the more grain refinement will be via the dynamic recrystallization at intermediate and high temperatures. Consequently, as the strain is increased to 0.45, the number of low-angle grain boundaries further increases while the twins nearly disappear, as seen in Figs. 8(g)-(i). It is obviously found that the number fraction of boundaries with misorientation angles less than 15° decreases while the number fraction of boundaries larger than 15° increases with the strain increasing from 0.45 to 0.60, as indicated in Figs. 8(j)-(l). It is suggested that the substructure with low-angle grain boundaries around initial grain boundaries absorbs subsequent dislocations and then develops into fine grains with high-angle boundary during compression, which is considered as a CDRX process [5,6,19,27-29]. The CDRX mechanism has been described by SITDIKOV and KAIBYSHEV [6], and VALLE et al [27].

In the initial stage of nucleation, the interaction between non-basal dislocation slips and cross slips leads to the formation of dislocation cells at both mantle regions of initial grains and twin boundaries. Then, the formation of sub-grain boundaries takes place by extended recovery. Ultimately, the high angle boundaries form via sub-boundary migration and coalescence. In order to further investigate the CDRX mechanism of the alloy during compression at 350 °C, the TEM images of the specimen are shown in Fig. 9. Plenty of pile-up and tangling dislocations around grain boundaries are observed. It is indicated that the dislocation cells transform from dislocations tangling, and then gradually develop into sub-grains at strain of 0.15 (as shown in Figs. 9(a) and (b)). As a result, serrated grain boundaries occur (as indicated in Figs. 4(b), 4(d) and Fig. 9(d)). The acceleration of DRX via dislocations increment results in the formation of the necklace-like structure around the large primary grains, which further confirms the activation of CDRX mechanism at 350 °C during compression.

Fig. 8 Inverse pole figures (a, d, g, j), grain boundary angle distribution figures (b, e, h, k) and twin distribution figures (c, f, i, l) of as-compressed alloy at 350 °C with different strains

3.4 Texture evolutions of Mg-Zn-Er alloy

The texture intensity as well as texture type shows a strong dependence on both temperature and strain. In particular, the texture intensity is sensitive to temperature. Figure 10 shows the pole figures of specimens after compression with different strains. When the temperature is 200 °C, the plastic deformation mainly relies on the basal slip and the pyramidal twin, which leads to a strong texture at a low strain owing to lower deformation temperature, as shown in Figs. 10(a) and (b). The experimental results show that the main texture consists of <0001> and <0001>. At the strain of 0.45, the main texture turns into , as indicated in Figs. 10(c) and (d). However, the texture intensity changes little, which is ascribed to the twin- dominant deformation mechanism. When the temperature raises to 350 °C, the texture feature of the specimen shows a significant difference, compared with that at 200 °C. It is observed that the texture component tilts about 45° from the CD at the strain of 0.08, as shown in Fig. 10(e). The maximum intensity value of the texture is ~59.49, which is far higher than that of the as-compressed specimen at 200 °C, i.e., 34.38. This implies that the decrease in twin intensity is not beneficial to modifying the texture. On the contrary, the twin has an obvious effect on weakening the texture in the initial deformation. The coarse grain structure produces a large number of twins in the process of deformation, which offers a possible means to modify the texture via rotating the basal plane by 86° [30]. With improvement of the strain, the texture component gradually shifts to  texture (Figs. 10(f) and (g)). Ultimately, the texture type is <0001>, as shown in Fig. 10(h). Meanwhile, the texture intensity drops rapidly as the strain increases, which is also ascribed to the CDRX mechanism [1].

Fig. 9 TEM images of as-compressed alloy at 350 °C with different strains

Fig. 10 Pole figures of specimens after compression with different strains at 200 °C (a-d) and 350 °C (e-h)

3.5 Precipitations of Mg-Zn-Er alloy during compression

Figure 11 shows the bright field TEM images of precipitates in the as-compressed alloy at the strain of 0.6. At 200 °C, there are large amounts of precipitates in both matrix and grain boundaries (Figs. 11(a) and (b)). The dimensions of particles at grain boundaries range 50-100 nm. Furthermore, the corresponding high- resolution TEM (HRTEM) and Fourier transformation (FFT) shows a typical nano-scale icosahedral quasicrystalline phase (I-phase, Mg₃Zn₆Er, as seen in Fig. 11(b)) with 3-fold symmetry [31]. Meanwhile, the needle or rod-shaped precipitates with several nanometers in width and 50 nm in length are dispersed in matrix, and they correspond to the Mg-Zn intermetallic phase [32]. However, it is found that the number of nano-scale I-phase (in 2-fold symmetry) in matrix increases at 350 °C, as shown in Figs. 11(c) and (d). Also, the nano-scale I-phases distribute more uniformly.

Previous investigations have shown that the thermo-mechanical treatment can not only destroy the large-size dendritic I-phase, but also promote the formation of a large amount of nano-scale I-phase, rendering dispersive distribution in matrix [31,32]. In this work, the deformation mechanisms of the alloy are main basal slip and twinning at 200 °C. Therefore, a considerable amount of deformation twin boundaries act as strong obstacles for dislocations motion, resulting in stress concentration, which leads to the precipitation of the nano-scale I-phase at grain boundaries. The stress plays an important role in introducing precipitations of the nano-scale I-phase and the Mg-Zn phase at 200 °C. When the deformation temperature is increased to 350 °C, the main precipitates are only the nano-scale I-phases. The nano-scale I-phases pin and hinder the dislocations motion, which leads to dislocation tangles and dislocation cells (Fig. 11(d)).

As stated above, the temperature is one of the key factors that influence the plastic deformation mechanism of Mg-Zn-Er alloy. The CRSS of non-basal slips and twinning are difficult to be activated compared with the basal slip at room temperature. The twinning and non-basal slips can be started and two competing deformation mechanisms are present when the high stress concentration is caused by inelastic strain. However, the twin nucleation is a process of stress-induced activation rather than heat-induced activation compared with the non-basal slips [29,33]. Therefore, the number of ＜a+c＞slip is large at room temperature and drops rapidly as the temperature is increased. Meanwhile, the twinning is expected to be relatively temperature-insensitive. Consequently, when the compression temperature is 200 °C, the twinning always exists in the whole process of the deformation and it is one of the main deformation mechanisms. With the strain of 0.6, the TDRX mechanism starts to happen because of the activation of the non-basal slip. However, the increase of temperature up to 350 °C leads to the activation of the non-basal slip, which results in the occurrence of the CDRX instead of the twins.

Fig. 11 Bright field TEM images of precipitates of as-compressed alloy with strain of 0.6 at different temperatures

4 Conclusions

1)  The temperature has a great influence on DRX mechanism. At the low temperature of 200 °C, the twinning is one of the main deformation mechanisms during the whole compression process. The TDRX begins to happen at the strain of 0.6 due to the activation of non-basal slips. Ultimately, the dominant features of microstructures are double twins and shear bands. At the high temperature of 350 °C, the twinning is still one of the main deformation mechanisms at a low strain. However, with the increase in deformation degree, CDRX substitutes the twining, and the final microstructure consists of necklace-like structures and initial grains.

2) The twin bands are mainly composed of dislocation walls, second twins, double twins, and micro-twins. These fine second twins are nuclei of TDRX. The shear begin began to form as the twin bands are consumed by ultrafine TDRX grains because of the activation of non-basal slips.

3) In the initial deformation stage, more tension twins result in the weaker texture. As the strain is increased, the increase in numbers of DRX grains noticeably decreases the texture at 350 °C. In addition, some kinds of nano-scale second phases precipitate during the deformation. The new precipitates of Mg-Zn binary phase and nano-sale I-phases are observed at 200 °C. The nano-scale I-phases are found as main precipitates during deformation at 350 °C.

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热压缩过程中温度对Mg-Zn-Er合金显微组织及织构演化的影响

刘金学¹，刘 轲¹，杜文博¹，李淑波¹，王朝辉¹，杜 宪¹，孙翠翠²

1. 北京工业大学 材料科学与工程学院，北京 100124；

2. 山东科学院，济南 250014

摘  要：通过光学显微镜、场发射扫描电镜和透射电镜研究热压缩过程中Mg-Zn-Er 合金的显微组织及织构的演化。结果表明，温度对动态再结晶(DRX)具有很大的影响。当温度为200 °C、应变量为0.6时，由于应力集中使得非基面滑移＜a+c＞位错被激活，孪生动态再结晶机制(TDRX)开始启动。当温度为350 °C时，围绕着初始晶粒的项链状结构出现，这是典型的连续动态再结晶机制(CDRX)。动态再结晶对弱化织构具有非常重要的影响，同时低温下孪生对弱化织构也起到一定的作用。研究还发现，当温度从200 °C提高到350 °C时，由于动态再结晶形核位点从孪晶界向初始晶界转移，织构减弱。

关键词：Mg-Zn-Er 合金；热压缩；动态再结晶；孪生；织构

(Edited by Bing YANG)

Foundation item: Project (2172013) supported by the Natural Science Foundation of Beijing, China; Project (2016YFB0301101) supported by the National Key Research and Development Program, China; Project (KZ201810005005) supported by Key Science and Technology Program of Beijing Municipal Commission of Education, China

Corresponding author: Ke LIU; Tel/Fax: +86-10-67392917; E-mail: lk@bjut.edu.cn

DOI: 10.1016/S1003-6326(18)64866-6