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Microstructure, texture and mechanical properties of as-extruded Mg–Zn–Er alloys
Materials Science & Engineering A 581 (2013) 31–38
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Microstructure, texture and mechanical properties of as-extruded Mg–Zn–Er alloys Qingfeng Wang, Wenbo Du n, Ke Liu, Zhaohui Wang, Shubo Li, Kai Wen College of Materials Science and Engineering, Beijing University of Technology, Beijing 100124, PR China
art ic l e i nf o
a b s t r a c t
Article history: Received 5 January 2013 Received in revised form 13 May 2013 Accepted 15 May 2013 Available online 23 May 2013
In the present work, the microstructure, texture and mechanical properties of the as-extruded Mg–xZn– yEr (x ¼3, 6 and 9; y¼ 0.5, 1 and 1.5; x/y¼6) alloys were investigated. The results showed that the I-phase in the as-cast alloys was destroyed and lots of nano-scale quasicrystalline particles were precipitated within the matrix during hot extrusion process. With the increase in the volume fraction of the primary I-phase, the particle-stimulated nucleation (PSN) of recrystallization was activated. It was suggested that the recrystallization via PSN should lead to the nucleation of new grains with a high orientation mismatch to the parent grains, making the texture gradually weaken. Significant strengthening is achieved in the Mg–9Zn–1.5Er alloy, of which the ultimate tensile strength (sb) and the yield tensile strength (s0.2) were 319 MPa and 192 MPa, respectively, with an elongation of 22.3%. & 2013 Elsevier B.V. All rights reserved.
Keywords: Mg–Zn–Er alloy Quasicrystal Microstructure Texture Mechanical properties Particle-stimulated nucleation
1. Introduction As the lightest metallic structural materials, magnesium alloys have been received a great attention in the last decade due to their good potential for use in automotive and aerospace applications [1]. However, the mechanical properties of the as-cast magnesium alloys are so inferior that they cannot satisfy the use in practice, especially for the Mg–Zn based alloys. In order to improve the mechanical properties of the alloys, the thermomechanical working process, such as hot extrusion (HE) and rolling, is usually carried out. A great enhancement in strength is usually achieved after thermomechanical working due to the refinement of microstructure (mainly attributed to the dynamic recrystallization) [2], the precipitation of second phase [3] and the relative strong texture [4]. However, the strong texture generally results in a plastic anisotropy, which may be responsible for the early failure of the alloys [5]. Eliminating plastic anisotropy via modifying texture has been pointed out by Hantzsche [5], Mukai [6] and Cottam [7], they all considered that grain refinement as well as modifying texture was an effective way to improve strength and elongation simultaneously [8]. Much work has been done on the modification of texture for magnesium alloys. For example, Kim et al. [9] reported that the equal channel angular process (ECAP) was an effective method to
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modify texture, and results in higher ductility. Mishra et al. [10] reported that a significant increase in elongation of pure magnesium was obtained with the addition of 0.2% cerium due to the change in texture of the extruded alloy in which the basal slip was favored. Besides, Sadeghi et al. [11] revealed that the random texture was associated with the particle stimulated nucleation (PSN) of recrystallization. Although the effects of PSN on promoting recrystallization were well reported for aluminum alloys [12,13], the recrystallization via PSN has also recently been observed and reported in magnesium alloys. For example, the Al–Sr particles in the Mg–3Al–1Zn–xSr (x¼ 0.4–0.8 wt%) alloys increased the potential nucleation of PSN, resulting in new grains with high orientation mismatch to the parent grains [11]. It may be considered that the PSN of recrystallization activated by thermally stable particles will be an effective way to modify the texture of magnesium alloys. The thermally stable icosahedra quasicrystalline phase (Mg3RE1Zn6, I-phase) has a low interfacial energy and results in a strong bonding to the I-phase/matrix interface [14]. Furthermore, the I-phase plays an important role in weakening the texture of Mg–Zn–Y alloys via DRX during hot rolling [15]. Both the strong bonding interface and random texture are beneficial to the good combination of elongation and strength for the wrought Mg–Zn–RE (rare earth) alloys [15–18]. However, up to now, the investigation on mechanical properties of magnesium alloys rarely focused on the influence of the I-phase on DRX via PSN. In the present work, the microstructure, texture and mechanical properties of the as-extruded Mg–xZn–yEr alloys (x ¼3, 6 and
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9; y ¼0.5, 1, and 1.5; x/y¼ 6, identified as alloys A, B and C, respectively) were investigated. The effect of the I-phase on the recrystallization via PSN and the effects of microstructure and texture on the mechanical properties of the as-extruded Mg–Zn–Er alloys were discussed.
2. Experimental procedures The as-cast alloys were prepared from the pure Mg (99.99 wt%), pure Zn (99.9 wt%) and Mg–30 wt% Er master alloys in a graphite crucible in an electric resistance furnace under an anti-oxidizing flux. The melt about 1200 g was poured into a steel mold. At last, an ingot with a size of 33 mm 120 mm 200 mm was obtained. The as-cast samples were annealed at 400 1C for 10 h, cooled down in 70 1C water, and then machined into ingots with a diameter of 35 mm. Then, the ingots were hot extruded into rods with a ratio of 10 at 250 1C, and the extrusion speed was 3 mm/s. The chemical compositions of alloys have been analyzed by an X-ray fluorescence (XRF) analyzer, as shown in Table 1. The phase analysis was performed by X-ray diffraction (XRD) with Cu Kα radiation. The microstructure observations were carried out by an optical microscope (OM, Zeiss-Imager A2m), a scanning electron microscope (SEM, HITACHI S-450) equipped with an energy dispersive spectroscope (EDS), a scanning electron microscope (SEM, JEOLJSM 6500F) equipped with Electron BackScatter Diffraction (EBSD), and a transmission electron microscope (TEM, JEMTable 1 Chemical compositions of Mg–Zn–Er alloys. Nominal alloy
Alloy A Alloy B Alloy C
Composition (wt%)
Zn/Er (wt%)
Zn
Er
Mg
3.0 5.9 8.9
0.5 1.0 1.4
Bal. Bal. Bal.
6.0 5.9 6.4
2000FX, JEOL). The samples for OM and SEM were mechanically polished and etched in a solution of 4 mL nitric acid and 96 mL ethanol. Specimens for TEM were prepared by electro-polishing and ion beam milling at an angle of incidence less than 101. Samples for texture analysis were prepared by electropolishing with a solution of 60% methanol, 30% glycerol and 10% nitric acid at 25 1C (voltage of 3–4 V) for 5–10 s. Tensile test was carried out by using a DNS-20 universal testing machine under a constant speed of 1.0 mm/min at room temperature. Specimens for the tensile test were made into dog-bone with a size of 5 mm gauge diameter and 25 mm gauge length. Three specimens were tested for each sample. All the tensile samples were sectioned in parallel to the extrusion direction.
3. Experimental results 3.1. Microstructure of the as-cast and as-extruded alloys Fig. 1a–c shows the SEM images of the as-cast alloys. It indicates that all of the as-cast alloys are composed of the interdendritic microstructure. The secondary phases with different morphologies can be found at both the matrix and interdendritic boundaries. In addition, the secondary phase at interdendritic boundaries becomes larger and more continuous and their volume fraction increases with the increase in the content of alloying elements. As reported in Refs. [19–22], the addition of RE into Mg–Zn based alloys would lead to the precipitation of I-phase on the basis of weight ratio of Zn/RE. Li [22] reported that the main secondary phase in the as-cast Mg–Zn–Er alloys was the I-phase when the Zn/Er weight ratio was ∼6. Fig. 1d shows the TEM image and corresponding SAED pattern of the alloy B. It indicates the distinct characteristic of the I-phase, which confirms that the main phase of the as-cast alloy in the present investigation is also the I-phase. Fig. 2 shows the OM images of the as-solution alloys A, B and C, annealed at 400 1C for 10 h. It indicates that the I-phase in the
Fig. 1. (a) SEM image of the as-cast alloy A, (b) SEM image of the as-cast alloy B, (c) SEM image of the as-cast alloy C and (d) TEM image of the I-phase and the corresponding SAED pattern with two-fold symmetry obtained from alloy B.
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alloys has not disappeared during the solution treatment. The existence of the I-phase suggests that it is thermally stable, which is attributed to the annealing temperature of 400 1C lowering than that of the phase transformation temperature of the I-phase (∼460 1C, reported in Ref. [22]). The I-phase in the assolution alloys can be classed into two types according to the morphologies, i.e. particle and anomalous. The I-phase mainly exists in the form of particle with a size of 2–5 μm in the alloy A. The I-phase with a particle morphology in the alloy B becomes more and bigger, and its size is about 3–10 μm. In addition, the
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coarse I-phase in an anomalous morphology appears in the alloy B with a width of 2–5 μm and length below 50 μm. The content of the coarse I-phase in anomalous morphology in the alloy C increases. Moreover, the coarse I-phase becomes continuous, while the content of particle I-phase reduces. The minimum particle diameter that could be reliably detected at the used magnification was ∼2 μm. Moreover, there are still a large number of coarse particles (about 10 μm diameter) left in the alloys. According to Ref. [23], the expected limit required for PSN is ∼1 μm. It implies that all of the I-phase particles in the three
Fig. 2. OM images of the as-solution alloys at 400 1C for 10 h: (a) alloy A, (b) alloy B and (c) alloy C.
Fig. 3. OM and SEM images of the as-extruded alloys: (a) OM image, (b) SEM image obtained from alloy A, (c) OM image, (d) SEM image obtained from alloy B, (e) OM image and (f) SEM image obtained from alloy C.
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Fig. 4. TEM bright field image of the as-extruded alloys: (a) the broken I-phase and the corresponding SAED pattern with five-fold symmetry obtained from the alloy B, (b) alloy A, (c) alloy B and (d) alloy C and the corresponding SAED patterns.
alloys have the tendency to activate recrystallization. In addition, the volume fractions of I-phase in alloy A, B and C, measured by an image analysis method, were ∼0.8%, 3.4% and 4.3%, respectively. Fig. 3 shows the OM images (a, c, and e) and SEM images (b, d, and f) of the as-extruded alloys along the extrusion direction (ED). The alloy A shows a bimodal grain structure consisting of fine grains and unDRXed region, whereas a fine microstructure with uniform grain size distribution is produced in the alloys B and C. It clearly suggests that the DRX took place during extrusion processing. Besides, the primary I-phase particles are destroyed during hot extrusion, distributing along the ED uniformly (see Fig. 3b, d and f), and the size of the broken I-phase is in a range of 0.5–5.0 mm. These broken I-phase particles may play an important role in limiting the growth of DRX grains [24]. In addition, it is found that the volume fraction of the broken particles with the size of ∼5 mm increases gradually as the volume fraction of the I-phase in the as-cast alloys increases. Fig. 4 shows the TEM images and corresponding SAED patterns of the microstructure obtained from the extruded alloys A, B and C. It can be obviously found that the broken interface between the I-phase and the matrix is completely healed by the thermally activated processes. Another interesting phenomenon is that many nano-scale particles can be found within the matrix of the asextruded alloys A, B and C. The nano-scale particles in elliptic or spherical morphology are also confirmed as the I-phase in terms of the SAED pattern with five-fold axes, as seen in the inset of Fig. 4d.
These nano-scale I-phase particles are precipitated during hot deformation, and they were also found in Mg–Zn–Y and Mg–Zn–Gd alloys [15,25–30]. Singh et al. [27] reported that the nano-scale I-phase precipitated during extrusion at 400 1C, and there was a definite orientation relationship between the nano-scale I-phase and the α-Mg matrix, i.e. sharply faceted on two-fold planes which were on the basal and prismatic planes of the α-Mg matrix. Kim et al. [28] reported that the nano-scale I-phase precipitated in the vicinity of the large eutectic I-phase during annealing at 420 1C for 48 h of the as-cast Mg96Zn3.4Y0.6 alloy (in at%). The precipitation of the nano-scale I-phase may be related with the defects introduced by the deformation, as reported by Liu [31].
3.2. Texture Fig. 5 shows the inverse pole figure (IPF) maps of the asextruded alloys A, B and C taken parallel to the extrusion direction obtained by using EBSD. In the maps, the red color stands for the (0002) basal plane, and the blue color represents the plane lying 901 away from the (0002) plane. The nearly identical color of two grains means that the misorientation between the grains is small. On the whole, the three as-extruded alloys are mainly composed of the grains exhibiting their (0002) basal plane. Moreover, the orientation of the grains becomes more dispersed as the content of the I-phase increases. Also, the as-extruded alloys include a nearly
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Fig. 5. Inverse pole figures map of the as-extruded alloys: (a) alloy A, (b) alloy B and (c) alloy C.
uniform grain structure which can be considered as a reasonably equiaxed grain. Fig. 6 displays the pole figures of the as-extruded alloys A, B and C. It indicates that the three alloys exhibit a strong (0002) basal texture. Specially, the alloy A has a very strong (0002) 〈101̄0〉 fiber texture which was found in the as-extruded Mg–Zn–Y–Zr alloys universally [24], with a maximum intensity value of ∼9.8. In addition, the (0002) basal plane of the most grains tends to be parallel to the extrusion direction (ED) [32]. The importance in Fig. 6 is that the increase in the volume fraction of the I-phase results in the transformation of texture. In the as-extruded alloys B and C, the fiber texture is destroyed and the orientation of the grains is more dispersed. The maximum intensity values of the texture in the as-extruded alloys B and C are ∼9.3 and 4.5, respectively. Moreover, with the content increment of the I-phase, the basal planes tend to tilt approximately 151 in the alloy B and 251 in the alloy C away from the ED. In order to clearly understand the micro-texture, the fraction of boundaries as a function of the misorientation angle for adjacent grains was measured. The results are listed in Table 2 and the normalized misorientation distribution histograms are shown in Fig. 7. It can be seen that there is a peak at ∼301 in the alloy A, which suggests that the strong basal texture is developed [33]. In addition, the fraction of the high angle grain boundaries (HAGBs) in the as-extruded alloys A, B and C is 89.3%, 91.1% and 92.8%,
respectively. As reported [34], the addition of RE elements leads to a random texture due to the DRX activated by the PSN. The recrystallized grain via PSN, which can provide an effective way to weaken the texture, is different from that produced by other recrystallization mechanisms [34]. In the present investigation, the precipitation of the I-phase leads to the formation of the recrystallized grains via PSN and the fraction of the HAGBs increases with the content of the I-phase particles. 3.3. Mechanical properties The stress–strain curves and the tensile properties of the extruded alloys at room temperature are shown in Fig. 8 and listed in Table 3, respectively. It can be seen that the UTS display a tendency to increase with the increase in the volume fraction of the I-phase. But the yield tensile strength (YTS) is converse on the whole. The YTS of the alloy A is about 207 MPa, whereas those of the alloys B and C are about 160 MPa and 192 MPa, respectively. Comparing with that of the alloys B and C, the YTS of the alloy A is greatly improved by 47 MPa and 15 MPa, respectively. And the elongation shows a trend of first rise after falling. Usually, the improvement of the strength is closely related with the grain size, texture and precipitation in wrought magnesium alloys [2,3]. Comparing with the results of grain size, texture and precipitation of the extruded alloys A, B and C, it is not difficult to
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Fig. 6. Pole figures of the as-extruded alloys, the RD standing for the ED: (a) alloy A, (b) alloy B and (c) alloy C.
Table 2 Fraction of boundaries with different misorientations in the as-extruded alloys A, B and C. Alloy
Alloy A Alloy B Alloy C
Number fraction (%) 2–151
15–1801
10.7 8.9 7.2
89.3 91.1 92.8
Average misorientation angles (deg)
51.28 52.95 51.74
understand the reasons for that the YTS of the alloy A is higher than those of the alloys B and C. They are mainly due to the finer grain size and stronger basal texture of the alloy A. The average grain size of the alloy A is ∼2.7 mm, while the average grain size of the alloy B is ∼4.7 mm. According to Hall–Petch relationship [35], the YTS is related inversely-proportional with the grain size. Besides, the maximal value of the texture of the alloy A is ∼9.8, which is the highest among the alloys. It is well known that a typical fiber texture parallel to the extrusion direction results in limiting the movement of the basal slip, because the Schmid factor was very small, which improves the YTS greatly [36]. Fig. 9 shows the fracture surfaces obtained from the alloys A, B and C after deformation at room temperature. The morphology of
the fracture surface exhibits a ductile fracture pattern, which is dominated by the void formation and combination, and lots of dimples can be found in the three alloys. It is also found that some cracked particles (∼5 mm) in the alloys B and C, especially in the alloy C. During deformation, the coarse particles were broken under the stress concentration and lead to an early failure.
4. Discussion Recently, it is found that the addition of RE elements and/or yttrium (RE/Y) elements leads to a weak texture because of the DRX. The mechanisms responsible for of the DRX usually include continuous dynamic recrystallization (CDRX), discontinuous dynamic recrystallization (DDRX), twin dynamic recrystallization (TDRX) and PSN. It is suggested that the DRX via PSN is related with the second phase particles [23]. During deformation of a particle-containing alloy, the enforced strain gradient in the vicinity of a non-deforming particle creates a region of high dislocation density and large orientation gradient, thus both high dislocation density and large orientation promote nucleation of recrystallization [23]. The random texture may also be ascribed to the DRX via PSN. Therefore, the texture in the alloys containing RE/Y is clearly different from the conventional basal texture of alloys [34]. For example, Tang et al. [37] already produced an alloy of
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Fig. 8. Stress–strain curves of the as-extruded alloys A, B and C tested under a constant speed of 1.0 mm/min at room temperature.
Table 3 Tensile properties of the as-extruded alloys A, B and C at room temperature.
Fig. 7. Normalized misorientation distribution histograms of the as-extruded alloys: (a) alloy A, (b) alloy B and (c) alloy C.
AZ31 by hot extrusion (370 1C, 1.4 m/min) with the maximum intensities of the (0002) plane is ∼45. In magnesium alloys, the effect of large particles on promoting recrystallization via PSN is well known. In order to examine whether the primary I-phase particles have the effect on promoting recrystallization via PSN in the present investigation, the hot extrusion process was interrupted when the material emerged ∼10 mm from the die and the alloy then quenched in water. The schematic view as well as the OM and TEM micrographs of the interrupted extrusion sample is shown in Fig. 10. The OM micrograph indicates that there are a great number of DRX grains around the coarse secondary particles. The DRX is suggested to be relative with the mechanism of the PSN, making the new formation of the DRX grains with great difference in orientation [11]. Fig. 10b shows the TEM observation of the sub-grains in the samples. It can be found that several DRX grains with size in the range about 1–3 mm formed in the vicinity of the broken I-phase. Besides, the grain boundary between the grains 1 and 2 is ∼141. The effect of large particles on promoting recrystallization via PSN is also observed in the other magnesium alloys. In the Mg–Mn alloys, the texture appears to be random due to the few grains which form via PSN [38]. In the Mg–3% Al–1% Zn–(0.4–0.8%) Sr alloys, the PSN was activated during hot extrusion at 350 1C, which
Alloy
UTS (MPa)
YTS (MPa)
E (%)
Alloy A Alloy B Alloy C
281 282 319
207 160 192
20.4 26.1 22.3
led to the nucleation of new grains with high orientation mismatch to the parent grains [11]. In the present investigation, the presence of the primary I-phase particles acted as nucleation sites for the PSN and made the texture to be random as the content of the I-phase particles increase. As a result, the maximum intensity values of the texture in the as-extruded alloys decreased with the increase in the content of the I-phase particles, which contributed to the plasticity of the as-extruded alloys. The elongation is affected by the texture intensity and the second phase in this study. The texture intensity decreases with the increase in content of the I-phase after hot extrusion, which results in the increase in the elongation. But the alloy C still contains some large particles with a size about 5 μm. During deformation, the coarse particles were broken under the stress concentration and lead to an early failure. Therefore, the elongations show a trend to rise first and then fall down.
5. Conclusions The effects of the I-phase on the microstructure, texture and mechanical properties of the as-extruded Mg–Zn–Er alloys were investigated. Some conclusions can be summarized as follows: (1) The I-phase could be broken and distributed along the extrusion direction, and the DRX occurred during hot extrusion. The primary I-phase particles worked as nucleation sites for the DRX via PSN. The higher volume fraction of the I-phase particles resulted in more extent of the DRX in the asextruded alloys. (2) The DRX via PSN contributed to the refined microstructure and random texture. As a result, the Mg–9Zn–1.5Er alloy exhibited good mechanical properties. The UTS (sb) and the YTS (s0.2) of the as-extruded alloy were 319 MPa and 192 MPa, respectively, with an elongation of 22.3%.
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Fig. 9. SEM fractographs of the as-extruded alloys: (a) alloy A, (b) alloy B, (c) alloy C and (d) image at a higher magnification of the as-extruded alloy C.
Fig. 10. Schematic map of the interrupted extrusion sample of Mg–9Zn–1.5Er alloy (a) the optical micrograph and (b) the TEM micrograph of DRX grains and I-phase with its SAED pattern.
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Microstructure, texture and mechanical properties of as-extruded Mg–Zn–Er alloys Qingfeng Wang, Wenbo Du n, Ke Liu, Zhaohui Wang, Shubo Li, Kai Wen College of Materials Science and Engineering, Beijing University of Technology, Beijing 100124, PR China
art ic l e i nf o
a b s t r a c t
Article history: Received 5 January 2013 Received in revised form 13 May 2013 Accepted 15 May 2013 Available online 23 May 2013
In the present work, the microstructure, texture and mechanical properties of the as-extruded Mg–xZn– yEr (x ¼3, 6 and 9; y¼ 0.5, 1 and 1.5; x/y¼6) alloys were investigated. The results showed that the I-phase in the as-cast alloys was destroyed and lots of nano-scale quasicrystalline particles were precipitated within the matrix during hot extrusion process. With the increase in the volume fraction of the primary I-phase, the particle-stimulated nucleation (PSN) of recrystallization was activated. It was suggested that the recrystallization via PSN should lead to the nucleation of new grains with a high orientation mismatch to the parent grains, making the texture gradually weaken. Significant strengthening is achieved in the Mg–9Zn–1.5Er alloy, of which the ultimate tensile strength (sb) and the yield tensile strength (s0.2) were 319 MPa and 192 MPa, respectively, with an elongation of 22.3%. & 2013 Elsevier B.V. All rights reserved.
Keywords: Mg–Zn–Er alloy Quasicrystal Microstructure Texture Mechanical properties Particle-stimulated nucleation
1. Introduction As the lightest metallic structural materials, magnesium alloys have been received a great attention in the last decade due to their good potential for use in automotive and aerospace applications [1]. However, the mechanical properties of the as-cast magnesium alloys are so inferior that they cannot satisfy the use in practice, especially for the Mg–Zn based alloys. In order to improve the mechanical properties of the alloys, the thermomechanical working process, such as hot extrusion (HE) and rolling, is usually carried out. A great enhancement in strength is usually achieved after thermomechanical working due to the refinement of microstructure (mainly attributed to the dynamic recrystallization) [2], the precipitation of second phase [3] and the relative strong texture [4]. However, the strong texture generally results in a plastic anisotropy, which may be responsible for the early failure of the alloys [5]. Eliminating plastic anisotropy via modifying texture has been pointed out by Hantzsche [5], Mukai [6] and Cottam [7], they all considered that grain refinement as well as modifying texture was an effective way to improve strength and elongation simultaneously [8]. Much work has been done on the modification of texture for magnesium alloys. For example, Kim et al. [9] reported that the equal channel angular process (ECAP) was an effective method to
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modify texture, and results in higher ductility. Mishra et al. [10] reported that a significant increase in elongation of pure magnesium was obtained with the addition of 0.2% cerium due to the change in texture of the extruded alloy in which the basal slip was favored. Besides, Sadeghi et al. [11] revealed that the random texture was associated with the particle stimulated nucleation (PSN) of recrystallization. Although the effects of PSN on promoting recrystallization were well reported for aluminum alloys [12,13], the recrystallization via PSN has also recently been observed and reported in magnesium alloys. For example, the Al–Sr particles in the Mg–3Al–1Zn–xSr (x¼ 0.4–0.8 wt%) alloys increased the potential nucleation of PSN, resulting in new grains with high orientation mismatch to the parent grains [11]. It may be considered that the PSN of recrystallization activated by thermally stable particles will be an effective way to modify the texture of magnesium alloys. The thermally stable icosahedra quasicrystalline phase (Mg3RE1Zn6, I-phase) has a low interfacial energy and results in a strong bonding to the I-phase/matrix interface [14]. Furthermore, the I-phase plays an important role in weakening the texture of Mg–Zn–Y alloys via DRX during hot rolling [15]. Both the strong bonding interface and random texture are beneficial to the good combination of elongation and strength for the wrought Mg–Zn–RE (rare earth) alloys [15–18]. However, up to now, the investigation on mechanical properties of magnesium alloys rarely focused on the influence of the I-phase on DRX via PSN. In the present work, the microstructure, texture and mechanical properties of the as-extruded Mg–xZn–yEr alloys (x ¼3, 6 and
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9; y ¼0.5, 1, and 1.5; x/y¼ 6, identified as alloys A, B and C, respectively) were investigated. The effect of the I-phase on the recrystallization via PSN and the effects of microstructure and texture on the mechanical properties of the as-extruded Mg–Zn–Er alloys were discussed.
2. Experimental procedures The as-cast alloys were prepared from the pure Mg (99.99 wt%), pure Zn (99.9 wt%) and Mg–30 wt% Er master alloys in a graphite crucible in an electric resistance furnace under an anti-oxidizing flux. The melt about 1200 g was poured into a steel mold. At last, an ingot with a size of 33 mm 120 mm 200 mm was obtained. The as-cast samples were annealed at 400 1C for 10 h, cooled down in 70 1C water, and then machined into ingots with a diameter of 35 mm. Then, the ingots were hot extruded into rods with a ratio of 10 at 250 1C, and the extrusion speed was 3 mm/s. The chemical compositions of alloys have been analyzed by an X-ray fluorescence (XRF) analyzer, as shown in Table 1. The phase analysis was performed by X-ray diffraction (XRD) with Cu Kα radiation. The microstructure observations were carried out by an optical microscope (OM, Zeiss-Imager A2m), a scanning electron microscope (SEM, HITACHI S-450) equipped with an energy dispersive spectroscope (EDS), a scanning electron microscope (SEM, JEOLJSM 6500F) equipped with Electron BackScatter Diffraction (EBSD), and a transmission electron microscope (TEM, JEMTable 1 Chemical compositions of Mg–Zn–Er alloys. Nominal alloy
Alloy A Alloy B Alloy C
Composition (wt%)
Zn/Er (wt%)
Zn
Er
Mg
3.0 5.9 8.9
0.5 1.0 1.4
Bal. Bal. Bal.
6.0 5.9 6.4
2000FX, JEOL). The samples for OM and SEM were mechanically polished and etched in a solution of 4 mL nitric acid and 96 mL ethanol. Specimens for TEM were prepared by electro-polishing and ion beam milling at an angle of incidence less than 101. Samples for texture analysis were prepared by electropolishing with a solution of 60% methanol, 30% glycerol and 10% nitric acid at 25 1C (voltage of 3–4 V) for 5–10 s. Tensile test was carried out by using a DNS-20 universal testing machine under a constant speed of 1.0 mm/min at room temperature. Specimens for the tensile test were made into dog-bone with a size of 5 mm gauge diameter and 25 mm gauge length. Three specimens were tested for each sample. All the tensile samples were sectioned in parallel to the extrusion direction.
3. Experimental results 3.1. Microstructure of the as-cast and as-extruded alloys Fig. 1a–c shows the SEM images of the as-cast alloys. It indicates that all of the as-cast alloys are composed of the interdendritic microstructure. The secondary phases with different morphologies can be found at both the matrix and interdendritic boundaries. In addition, the secondary phase at interdendritic boundaries becomes larger and more continuous and their volume fraction increases with the increase in the content of alloying elements. As reported in Refs. [19–22], the addition of RE into Mg–Zn based alloys would lead to the precipitation of I-phase on the basis of weight ratio of Zn/RE. Li [22] reported that the main secondary phase in the as-cast Mg–Zn–Er alloys was the I-phase when the Zn/Er weight ratio was ∼6. Fig. 1d shows the TEM image and corresponding SAED pattern of the alloy B. It indicates the distinct characteristic of the I-phase, which confirms that the main phase of the as-cast alloy in the present investigation is also the I-phase. Fig. 2 shows the OM images of the as-solution alloys A, B and C, annealed at 400 1C for 10 h. It indicates that the I-phase in the
Fig. 1. (a) SEM image of the as-cast alloy A, (b) SEM image of the as-cast alloy B, (c) SEM image of the as-cast alloy C and (d) TEM image of the I-phase and the corresponding SAED pattern with two-fold symmetry obtained from alloy B.
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alloys has not disappeared during the solution treatment. The existence of the I-phase suggests that it is thermally stable, which is attributed to the annealing temperature of 400 1C lowering than that of the phase transformation temperature of the I-phase (∼460 1C, reported in Ref. [22]). The I-phase in the assolution alloys can be classed into two types according to the morphologies, i.e. particle and anomalous. The I-phase mainly exists in the form of particle with a size of 2–5 μm in the alloy A. The I-phase with a particle morphology in the alloy B becomes more and bigger, and its size is about 3–10 μm. In addition, the
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coarse I-phase in an anomalous morphology appears in the alloy B with a width of 2–5 μm and length below 50 μm. The content of the coarse I-phase in anomalous morphology in the alloy C increases. Moreover, the coarse I-phase becomes continuous, while the content of particle I-phase reduces. The minimum particle diameter that could be reliably detected at the used magnification was ∼2 μm. Moreover, there are still a large number of coarse particles (about 10 μm diameter) left in the alloys. According to Ref. [23], the expected limit required for PSN is ∼1 μm. It implies that all of the I-phase particles in the three
Fig. 2. OM images of the as-solution alloys at 400 1C for 10 h: (a) alloy A, (b) alloy B and (c) alloy C.
Fig. 3. OM and SEM images of the as-extruded alloys: (a) OM image, (b) SEM image obtained from alloy A, (c) OM image, (d) SEM image obtained from alloy B, (e) OM image and (f) SEM image obtained from alloy C.
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Fig. 4. TEM bright field image of the as-extruded alloys: (a) the broken I-phase and the corresponding SAED pattern with five-fold symmetry obtained from the alloy B, (b) alloy A, (c) alloy B and (d) alloy C and the corresponding SAED patterns.
alloys have the tendency to activate recrystallization. In addition, the volume fractions of I-phase in alloy A, B and C, measured by an image analysis method, were ∼0.8%, 3.4% and 4.3%, respectively. Fig. 3 shows the OM images (a, c, and e) and SEM images (b, d, and f) of the as-extruded alloys along the extrusion direction (ED). The alloy A shows a bimodal grain structure consisting of fine grains and unDRXed region, whereas a fine microstructure with uniform grain size distribution is produced in the alloys B and C. It clearly suggests that the DRX took place during extrusion processing. Besides, the primary I-phase particles are destroyed during hot extrusion, distributing along the ED uniformly (see Fig. 3b, d and f), and the size of the broken I-phase is in a range of 0.5–5.0 mm. These broken I-phase particles may play an important role in limiting the growth of DRX grains [24]. In addition, it is found that the volume fraction of the broken particles with the size of ∼5 mm increases gradually as the volume fraction of the I-phase in the as-cast alloys increases. Fig. 4 shows the TEM images and corresponding SAED patterns of the microstructure obtained from the extruded alloys A, B and C. It can be obviously found that the broken interface between the I-phase and the matrix is completely healed by the thermally activated processes. Another interesting phenomenon is that many nano-scale particles can be found within the matrix of the asextruded alloys A, B and C. The nano-scale particles in elliptic or spherical morphology are also confirmed as the I-phase in terms of the SAED pattern with five-fold axes, as seen in the inset of Fig. 4d.
These nano-scale I-phase particles are precipitated during hot deformation, and they were also found in Mg–Zn–Y and Mg–Zn–Gd alloys [15,25–30]. Singh et al. [27] reported that the nano-scale I-phase precipitated during extrusion at 400 1C, and there was a definite orientation relationship between the nano-scale I-phase and the α-Mg matrix, i.e. sharply faceted on two-fold planes which were on the basal and prismatic planes of the α-Mg matrix. Kim et al. [28] reported that the nano-scale I-phase precipitated in the vicinity of the large eutectic I-phase during annealing at 420 1C for 48 h of the as-cast Mg96Zn3.4Y0.6 alloy (in at%). The precipitation of the nano-scale I-phase may be related with the defects introduced by the deformation, as reported by Liu [31].
3.2. Texture Fig. 5 shows the inverse pole figure (IPF) maps of the asextruded alloys A, B and C taken parallel to the extrusion direction obtained by using EBSD. In the maps, the red color stands for the (0002) basal plane, and the blue color represents the plane lying 901 away from the (0002) plane. The nearly identical color of two grains means that the misorientation between the grains is small. On the whole, the three as-extruded alloys are mainly composed of the grains exhibiting their (0002) basal plane. Moreover, the orientation of the grains becomes more dispersed as the content of the I-phase increases. Also, the as-extruded alloys include a nearly
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Fig. 5. Inverse pole figures map of the as-extruded alloys: (a) alloy A, (b) alloy B and (c) alloy C.
uniform grain structure which can be considered as a reasonably equiaxed grain. Fig. 6 displays the pole figures of the as-extruded alloys A, B and C. It indicates that the three alloys exhibit a strong (0002) basal texture. Specially, the alloy A has a very strong (0002) 〈101̄0〉 fiber texture which was found in the as-extruded Mg–Zn–Y–Zr alloys universally [24], with a maximum intensity value of ∼9.8. In addition, the (0002) basal plane of the most grains tends to be parallel to the extrusion direction (ED) [32]. The importance in Fig. 6 is that the increase in the volume fraction of the I-phase results in the transformation of texture. In the as-extruded alloys B and C, the fiber texture is destroyed and the orientation of the grains is more dispersed. The maximum intensity values of the texture in the as-extruded alloys B and C are ∼9.3 and 4.5, respectively. Moreover, with the content increment of the I-phase, the basal planes tend to tilt approximately 151 in the alloy B and 251 in the alloy C away from the ED. In order to clearly understand the micro-texture, the fraction of boundaries as a function of the misorientation angle for adjacent grains was measured. The results are listed in Table 2 and the normalized misorientation distribution histograms are shown in Fig. 7. It can be seen that there is a peak at ∼301 in the alloy A, which suggests that the strong basal texture is developed [33]. In addition, the fraction of the high angle grain boundaries (HAGBs) in the as-extruded alloys A, B and C is 89.3%, 91.1% and 92.8%,
respectively. As reported [34], the addition of RE elements leads to a random texture due to the DRX activated by the PSN. The recrystallized grain via PSN, which can provide an effective way to weaken the texture, is different from that produced by other recrystallization mechanisms [34]. In the present investigation, the precipitation of the I-phase leads to the formation of the recrystallized grains via PSN and the fraction of the HAGBs increases with the content of the I-phase particles. 3.3. Mechanical properties The stress–strain curves and the tensile properties of the extruded alloys at room temperature are shown in Fig. 8 and listed in Table 3, respectively. It can be seen that the UTS display a tendency to increase with the increase in the volume fraction of the I-phase. But the yield tensile strength (YTS) is converse on the whole. The YTS of the alloy A is about 207 MPa, whereas those of the alloys B and C are about 160 MPa and 192 MPa, respectively. Comparing with that of the alloys B and C, the YTS of the alloy A is greatly improved by 47 MPa and 15 MPa, respectively. And the elongation shows a trend of first rise after falling. Usually, the improvement of the strength is closely related with the grain size, texture and precipitation in wrought magnesium alloys [2,3]. Comparing with the results of grain size, texture and precipitation of the extruded alloys A, B and C, it is not difficult to
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Fig. 6. Pole figures of the as-extruded alloys, the RD standing for the ED: (a) alloy A, (b) alloy B and (c) alloy C.
Table 2 Fraction of boundaries with different misorientations in the as-extruded alloys A, B and C. Alloy
Alloy A Alloy B Alloy C
Number fraction (%) 2–151
15–1801
10.7 8.9 7.2
89.3 91.1 92.8
Average misorientation angles (deg)
51.28 52.95 51.74
understand the reasons for that the YTS of the alloy A is higher than those of the alloys B and C. They are mainly due to the finer grain size and stronger basal texture of the alloy A. The average grain size of the alloy A is ∼2.7 mm, while the average grain size of the alloy B is ∼4.7 mm. According to Hall–Petch relationship [35], the YTS is related inversely-proportional with the grain size. Besides, the maximal value of the texture of the alloy A is ∼9.8, which is the highest among the alloys. It is well known that a typical fiber texture parallel to the extrusion direction results in limiting the movement of the basal slip, because the Schmid factor was very small, which improves the YTS greatly [36]. Fig. 9 shows the fracture surfaces obtained from the alloys A, B and C after deformation at room temperature. The morphology of
the fracture surface exhibits a ductile fracture pattern, which is dominated by the void formation and combination, and lots of dimples can be found in the three alloys. It is also found that some cracked particles (∼5 mm) in the alloys B and C, especially in the alloy C. During deformation, the coarse particles were broken under the stress concentration and lead to an early failure.
4. Discussion Recently, it is found that the addition of RE elements and/or yttrium (RE/Y) elements leads to a weak texture because of the DRX. The mechanisms responsible for of the DRX usually include continuous dynamic recrystallization (CDRX), discontinuous dynamic recrystallization (DDRX), twin dynamic recrystallization (TDRX) and PSN. It is suggested that the DRX via PSN is related with the second phase particles [23]. During deformation of a particle-containing alloy, the enforced strain gradient in the vicinity of a non-deforming particle creates a region of high dislocation density and large orientation gradient, thus both high dislocation density and large orientation promote nucleation of recrystallization [23]. The random texture may also be ascribed to the DRX via PSN. Therefore, the texture in the alloys containing RE/Y is clearly different from the conventional basal texture of alloys [34]. For example, Tang et al. [37] already produced an alloy of
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Fig. 8. Stress–strain curves of the as-extruded alloys A, B and C tested under a constant speed of 1.0 mm/min at room temperature.
Table 3 Tensile properties of the as-extruded alloys A, B and C at room temperature.
Fig. 7. Normalized misorientation distribution histograms of the as-extruded alloys: (a) alloy A, (b) alloy B and (c) alloy C.
AZ31 by hot extrusion (370 1C, 1.4 m/min) with the maximum intensities of the (0002) plane is ∼45. In magnesium alloys, the effect of large particles on promoting recrystallization via PSN is well known. In order to examine whether the primary I-phase particles have the effect on promoting recrystallization via PSN in the present investigation, the hot extrusion process was interrupted when the material emerged ∼10 mm from the die and the alloy then quenched in water. The schematic view as well as the OM and TEM micrographs of the interrupted extrusion sample is shown in Fig. 10. The OM micrograph indicates that there are a great number of DRX grains around the coarse secondary particles. The DRX is suggested to be relative with the mechanism of the PSN, making the new formation of the DRX grains with great difference in orientation [11]. Fig. 10b shows the TEM observation of the sub-grains in the samples. It can be found that several DRX grains with size in the range about 1–3 mm formed in the vicinity of the broken I-phase. Besides, the grain boundary between the grains 1 and 2 is ∼141. The effect of large particles on promoting recrystallization via PSN is also observed in the other magnesium alloys. In the Mg–Mn alloys, the texture appears to be random due to the few grains which form via PSN [38]. In the Mg–3% Al–1% Zn–(0.4–0.8%) Sr alloys, the PSN was activated during hot extrusion at 350 1C, which
Alloy
UTS (MPa)
YTS (MPa)
E (%)
Alloy A Alloy B Alloy C
281 282 319
207 160 192
20.4 26.1 22.3
led to the nucleation of new grains with high orientation mismatch to the parent grains [11]. In the present investigation, the presence of the primary I-phase particles acted as nucleation sites for the PSN and made the texture to be random as the content of the I-phase particles increase. As a result, the maximum intensity values of the texture in the as-extruded alloys decreased with the increase in the content of the I-phase particles, which contributed to the plasticity of the as-extruded alloys. The elongation is affected by the texture intensity and the second phase in this study. The texture intensity decreases with the increase in content of the I-phase after hot extrusion, which results in the increase in the elongation. But the alloy C still contains some large particles with a size about 5 μm. During deformation, the coarse particles were broken under the stress concentration and lead to an early failure. Therefore, the elongations show a trend to rise first and then fall down.
5. Conclusions The effects of the I-phase on the microstructure, texture and mechanical properties of the as-extruded Mg–Zn–Er alloys were investigated. Some conclusions can be summarized as follows: (1) The I-phase could be broken and distributed along the extrusion direction, and the DRX occurred during hot extrusion. The primary I-phase particles worked as nucleation sites for the DRX via PSN. The higher volume fraction of the I-phase particles resulted in more extent of the DRX in the asextruded alloys. (2) The DRX via PSN contributed to the refined microstructure and random texture. As a result, the Mg–9Zn–1.5Er alloy exhibited good mechanical properties. The UTS (sb) and the YTS (s0.2) of the as-extruded alloy were 319 MPa and 192 MPa, respectively, with an elongation of 22.3%.
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Fig. 9. SEM fractographs of the as-extruded alloys: (a) alloy A, (b) alloy B, (c) alloy C and (d) image at a higher magnification of the as-extruded alloy C.
Fig. 10. Schematic map of the interrupted extrusion sample of Mg–9Zn–1.5Er alloy (a) the optical micrograph and (b) the TEM micrograph of DRX grains and I-phase with its SAED pattern.
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