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Effect of temperature on deformation mechanisms of the Mg88Co5Y7 alloy during hot compression
Materials Characterization 151 (2019) 553–562
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Effect of temperature on deformation mechanisms of the Mg88Co5Y7 alloy during hot compression ⁎
Z.Z. Penga, Q.Q. Jina, X.H. Shaoa, , Y.T. Zhoua, S.J. Zhenga, B. Zhanga, X.L. Maa,b, a b
T
⁎
Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, 110016 Shenyang, China School of Materials Science and Engineering, Lanzhou University of Technology, 287 Langongping Road, 730050 Lanzhou, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Magnesium alloy LPSO structure Deformation twin Recrystallization Intermetallics
The deformation mechanisms and microstructure evolution of Mg88Co5Y7 (at.%) alloy with long period stacking ordered (LPSO) phase during compression at 373–673 K have been studied using transmission electron microscopy. The alloy consists of α-Mg matrix, interdendritic LPSO phase, and other Mg-Co-Y intermetallic compounds. Increasing temperature leads to a general decrease in strength of the Mg88Co5Y7 alloy. A large number of {1012 } tension twins and deformation bands are activated in the matrix, whereas basal slip and deformation kink dominate in the LPSO structures during compression at 373 and 473 K. Non-basal slip in matrix and almost no dynamic recrystallization are responsible for high strength and the good ductility of sample deformed at 573 K. Recrystallization of Mg matrix occurs upon deformation at 673 K, dramatically lowering the corresponding strength. Further, the grain boundary pinning effect from LPSO phase, MgYCo4 phase, and broken Mg3(Co, Y) segments is supposed to account for the relatively high strength of Mg88Co5Y7 alloy at high temperatures.
1. Introduction Because of low strength and poor plasticity, commercial magnesium alloys as structural materials have limited application in the automobile industry and other industries. Mg-TM-RE (TM, transition metals; RE, rare earth elements) ternary alloys containing long period stacking ordered (LPSO) phases have received considerable attention during the past decades due to their excellent performance both at room temperature and elevated temperatures [1–8]. Of the ternary alloy systems that LPSO phases have been observed, Mg-TM-RE (TM = Ni, Cu and Zn; RE = Y, Gd, Tb, Dy, Ho, and Er) alloys [2–4,9] were reported to contain LPSO phases with AB′C′A stacking faults, where B′ and C′ layers are commonly enriched by TM/RE atoms [10]. The mechanical properties of these alloys consisting of Mg and LPSO phases have been extensively studied [1–7]. The deformation mechanisms include deformation kink [7], deformation twin and its interaction with LPSO phases [11,12] and non-basal dislocations in matrix [8,13]. Very recently, new polytypes of LPSO structures featuring AB′C stacking faults (B′ layer is enriched in Co/Y atoms), as well as some other intermetallic compounds, such as Mg24Y5, Mg3(Co, Y) and MgYCo4 phases, were characterized in Mg-Co-Y as-cast alloys [14–16]. The microstructural details (including morphology, crystal structure and volume fraction of the constituent phase) and phase transformation of the as-cast and annealed Mg88Co5Y7 (at.%) alloys ⁎
have been elaborated [17,18]. However, although the hardness measurements demonstrate that the LPSO structures in Mg-Co-Y systems are harder than those in Mg-Zn-Y alloys [7,14], there is very little information considering mechanical properties and related deformation mechanisms of Mg-Co-Y alloys containing AB′C-LPSO structure and other intermetallic compounds. In the present paper, we investigated the deformation behaviors of the as-cast Mg88Co5Y7 (at.%) alloy during hot compression. The microstructure evolution of the samples at different temperatures was characterized by scanning electron microscope (SEM) and transmission electron microscope (TEM). The corresponding mechanisms are also discussed. Further, we try to reveal the effects of the LPSO phase and intermetallic on the strength and ductility of the Mg88Co5Y7 alloy, which will shed some new light on optimizing the mechanical properties of the magnesium alloys by introducing distinctive microstructures. 2. Experimental procedures A ternary alloy with a nominal composition Mg88Co5Y7 (at.%) was prepared by melting the high purity Mg, Co and Mg-30wt%Y master alloys in a graphite crucible under protection of ultrahigh purity argon atmosphere in a high-frequency induction melting furnace, and then cooled down to the room temperature under the argon atmosphere in the furnace. Specimens with the dimensions 4 × 4 × 8 mm3 were cut
Corresponding authors. E-mail addresses: [email protected] (X.H. Shao), [email protected] (X.L. Ma).
https://doi.org/10.1016/j.matchar.2019.03.049 Received 10 December 2018; Received in revised form 9 March 2019; Accepted 31 March 2019 Available online 01 April 2019 1044-5803/ © 2019 Elsevier Inc. All rights reserved.
Materials Characterization 151 (2019) 553–562
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from the ingot by low-speed diamond cutting machine cooling by water. Compression experiment was carried out at 373 K, 473 K, 573 K, and 673 K and a strain rate of 1.0 × 10−3 s−1 in a Gleeble-3800 thermal simulation machine. We refer these samples S373, S473, S573, and S673 hereafter, respectively. Prior to being compressed the specimens were conductively heated to the setting temperatures at a heating rate of 5 K s−1 and held for 180 s for equilibration. The compression direction was parallel to the long axis of the specimens. The specimens were deformed to failure (S373 and S473) or stopped at the setting strain (80% for S573 and S673), then immediately water quenched to room temperature for microstructure analysis. The microstructures of specimens were examined by scanning electron microscopy (SEM) and scanning transmission electron microscopy (STEM). Thin foil samples for TEM were prepared by the conventional ion milling method, a Gatan precision ion polishing system (PIPS 691) with a liquid-nitrogen-cooled stage to avoid preferential thinning effects. The selected area electron diffraction (SAED) patterns and bright-field (BF) TEM images were obtained from the JEOL JEM 2100F FEG-TEM operating at 200 kV. The Titan3™ G2 60-300, which was operated at 300 kV and equipped with an energy-dispersive X-ray (EDX) detector and a high-angle annular dark-field (HAADF) detector, was used to obtain high-resolution STEM images and to measure chemical compositions.
Table 1 Compressive stress, yield stress and a compression ratio of Mg88Co5Y7 alloy at 373 K, 473 K, 573 K, and 673 K respectively. Compression stress σ (MPa)
Yield stress σ0.2 (MPa)
Compression ratio δ (%)
311 311 234 108
145 145 100 90
22 22 > 80 > 80
373 K 473 K 573 K 673 K
Table 2 Compressive stresses of Mg88Co5Y7 (at.%) alloys and other LPSO containing magnesium alloys. (The compressive stresses here are true stresses from the true stress-strain curves and the strain rate are all 1.0 × 10−3 s−1.) Magnesium alloys
Mg97Zn1Y2 (at.%) MgZn2Zr0.3Y5.8 (wt%) 693 K, 12 h MgZn2Zr0.3Y5.8 (wt%) 773 K, 20 h Mg-Gd-Y-Zn-Zr Mg-Zn-Mn-Y Mg88Co5Y7 (at.%)
Compression stress σ (MPa)
Ref.
373 K
473 K
573 K
673 K
223
217
311
317 311
190 164 210 213 246 234
96 75 89 50 79 108
[19] [20] [21] [22] [23] Present
This highlights that the Mg88Co5Y7 (at.%) alloy shows high compressive stress at high temperatures.
3. Results 3.1. True stress–true strain curves
other magnesium alloys listed in Table 2, we can conclude that the compressive strength of Mg88Co5Y7 alloy is superior to most cast [20] and homogenized LPSO containing magnesium alloys [21,22], nearly equivalent to the mechanical strength of extruded LPSO containing magnesium alloys [23,24], especially at 573 K and 673 K.
The compression stress–strain curves of the Mg88Co5Y7 alloy under a strain rate of 1.0 × 10−3 s−1 at 373 K, 473 K, 573 K, and 673 K are shown in Fig. 1. The curves for S373 and S473 are nearly coincident, with the peak stress of about 311 MPa and the compression ratio of 22%, suggesting that the deformation behavior at 373 K and 473 K should be similar. The strength for S573 decreases by 24.8% (from 311 MPa to 234 MPa), whereas the compression ratio improves over 264% (from 22% to above 80%). In contrast, the strength of S673 dramatically decreases to 108 MPa. The serrated flow may suggest dynamic strain aging due to the mutual interaction between dislocations and solute atoms during hot compression since solute atoms in Mg alloys containing LPSO structures tend to segregate along dislocations and stacking faults [19]. But the underlying mechanism would be investigated in the near future. Table 1 summarizes the compressive strength, yield strength (σ0.2) and a compression ratio of the Mg88Co5Y7 samples compressed at different temperatures. Compared with the compression properties of
3.2. Microstructural evolution and deformation mechanisms of Mg88Co5Y7 deformed alloy 3.2.1. Microstructural morphologies of Mg88Co5Y7 alloy deformed at different temperatures Fig. 2 are the back-scattering electron (BSE) SEM image of the Mg88Co5Y7 alloy, clearly presenting the evolution of intermetallic compounds during compression. As indicated by the small arrows in Fig. 2a, four kinds of intermetallics, LPSO phase, Mg24Y5 phase (about 0.9% volume fraction), Mg3(Co, Y) phase and MgYCo4 phase, coexist at the as-cast state. The details about their composition, crystal structure and volume fraction in the cast alloy have been illustrated in refs. [14, 17, 18]. Fig. 2b–d presents the deformation microstructures of S473, S573, and S673, respectively. Fig. 2b shows that the morphology of S473 is almost the same as that at as-cast state, except some of the LPSO phases are slightly curved. The S373 is structurally analogous to S473, which is not shown. For S573, Fig. 2c indicates that the intergrown LPSO/Mg3(Co, Y) phase is obviously curved, which may infer the formation of deformation kink. Fig. 2d shows that Mg matrix with LPSO phase and Mg3(Co, Y) phase at the grain boundary were elongated perpendicular to the loading axis during 673 K. Moreover, the Mg3(Co, Y) phase evolved into small particles and dispersed along the edges of the LPSO phases, as marked by the yellow arrows in the inserted zoomin image in Fig. 2d. MgYCo4 phase in Mg matrix seems to be undeformed during compression. In order to further illustrate the corresponding deformation mechanisms of Mg88Co5Y7 alloy at various temperatures, we studied the samples by TEM observation. Note that in this work the observation plane of the TEM sample is parallel to the compression axis.
Fig. 1. True stress–true strain curves obtained from compressive tests at 373 K, 473 K, 573 K, and 673 K under a constant strain rate of 1.0 × 10−3 s−1 respectively.
3.2.2. Deformation behavior at 473 K Profuse deformation twins and kinks occur in Mg matrix of S473. Each of the five bands in Fig. 3a, numbered “1–5”, exhibits twin 554
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Fig. 2. Back-scattering electron (BSE) images of the Mg88Co5Y7 alloys. (a) As-cast (The arrows in the enlarged area show the segregation regions of heavy elements at the interface between LPSO phase and Mg matrix); (b) S473; (c) S573 and (d) S673, showing the evolution of the microstructures. (For interpretation of the references to color in this figure, the reader is referred to the web version of this article.)
between {1012} twin and ISF here is consistent with that in compressed Mg97Zn1Y2 alloys [11]. Meanwhile, Co and/or Y atoms periodically segregated along the CTB a little bit farther from the intersection between twin and ISF, as shown in Fig. 4d. Additionally, Fig. 4 also demonstrated that heavy solute atoms are easily segregated to stacking faults and twin boundaries in this alloy. Basal dislocations, low and high angle grain boundaries (GBs) were produced in the LPSO phases. Fig. 5a is the BF-TEM image of one deformed LPSO phase, indicating large amounts of dislocations are activated during the deformation process. The dark line contrast along basal plane in the BF-TEM image acquired along the 〈1120〉 orientation of the LPSO phase (Fig. 5b), suggesting that < a > dislocations or basal SFs glide on basal planes, as indicated by the arrows. The accumulation of the defects described above can make the LPSO slightly bent and even result in deformation kink [8]. For example, multiple kinks in the LPSO phase were detected when the local stress increased at the interface between the 15R-LPSO and a harder intermetallic Mg3(Co, Y) phase, as shown in Fig. 5c. A very small amount of Mg3(Co, Y) phase was broken. Fig. 6a is a low magnification HAADF image of the intermetallic Mg3(Co, Y) in S473, showing the cracks propagate along the {1101} planes of the Mg3(Co, Y) phase. The Co2Y9 clusters in Mg3 (Co, Y) cell in the same row deflected from each other along the {1011} plane, as shown in Fig. 6b, where the fast Fourier transformation (FFT) image is inserted. Combining the crystal structure, we assume that the fracture plane of Mg3(Co, Y) is closely related to the chemical bonding of the atoms, which need further related simulation.
orientation with its neighboring band. The SAED pattern of the circle area in Fig. 3a indicates that the “1” and its neighbors show {1012} tensile twin orientation relationship, demonstrated by electron diffraction (ED) pattern in Fig. 3b. Actually, {1012} tensile twin is the main twin in the deformed Mg matrix under this deformation condition in our study. Fig. 3c and d is the BF-TEM and HAADF images of the deformation bands in the Mg matrix. They always show sharp interfaces with the Mg matrix, similar to the morphology of twins. Nevertheless, the deformation band always deviates from the zone axis when the matrix is rotated along 〈1120〉 or 〈1010〉 direction. This can be further proved by the strong contrast differences between the deformation band and the matrix in the BF-TEM and HAADF images. These deformation bands might be other twin bands rather than traditional {1012}, {1011} or {1013} twin bands [25], which will be explored in the near future. It should be mentioned the multiple twins and deformation bands are always generated together, crossing or paralleling to each other, as shown in Fig. 3a and c. The migrating of the twin boundary (TB) and its interaction with other defects are also frequently detected in S473. Fig. 4a and b shows HAADF images of the {1012} twin and its interaction with the intrinsic stacking fault (ISF), respectively. Fig. 4c is the corresponding high-resolution STEM image of the framed area in Fig. 4b. The TB consists of BP/PB boundary (B represents basal plane, P represents prismatic plane) and coherent twin boundary (CTB), respectively represented by solid and dashed lines. The stacking sequence of the ISF is “…ABAB′CBC…”, where the B′ layer is brighter due to a higher concentration of Co and/or Y elements compared with the adjacent magnesium layers. The stacking structure of AB′C is same as the stacking block of 15R, 12H and 21R-LPSO in the Mg-Co-Y alloy [14]. After twinning, the original AB′C fault structure converts to the ABAB stacking sequence in the twin, as shown in Fig. 4c. The interaction
3.2.3. Deformation behavior at 573 K Profuse dislocations in the Mg matrix were detected in S573, as shown in Fig. 7a. The pyramidal dislocation movement has more 555
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Fig. 3. Twins and deformation bands formed in the deformed Mg matrix. (a) and (b) Morphology of the twins and selected area electron diffraction (SAED) pattern of the {1012} twin; (c) and (d) BF-TEM and HAADF images of the deformation band, respectively.
as shown in Fig. 8.
advantages than twinning motivation in energy when the deformation temperature is above 200 °C [26]. Fig. 7b and c is the high-resolution TEM images of the formed boundary at location ① and ②, respectively, suggesting that the boundary parallel or vertical to the Mg basal plane is composed of the dissociated < c + a > dislocations and an array of < a > dislocations. Fig. 7d is the FFT pattern of Fig. 7c, indicating the small rotation angle of the basal planes is 3.6°. In contrast, deformation twins or kink bands were rarely observed in the matrix. The most prominent feature in S573 is the precipitation of nanoparticles in the Mg matrix nearby the LPSO phase, as shown in Fig. 8. Fig. 8a and b shows BF-TEM and corresponding HAADF image of the nanoparticles, respectively, indicating that the particles with polyhedron shape are in the range of 100–300 nm. To clarify the nanoparticles, the chemical composition and crystal structure of the particle were analyzed. Fig. 9 shows morphology and corresponding EDS mapping images of the nanoparticles, indicating that the nanoparticles are enriched in Y element and no detectable Co element. The specific chemical composition of the nanoparticle is estimated to be 88.99 ± 0.95 at.% Mg – 0.04 ± 0.02 at.% Co – 10.96 ± 0.97 at.% Y, which suggests that it should be MgeY binary phase. Fig. 10a–c is highresolution STEM images of the nanoparticles observing from various zone axis, and the corresponding FFT images of Fig. 10a–c are shown in Fig. 10d–f, respectively. Combined the comprehensive analysis of chemical composition and crystal structure, it's fair to say that the nanoparticle is Mg24Y5 phase. The atomic models of Mg24Y5 phase are superimposed in Fig. 10a–c, where blue and orange dots represent Y and Mg atoms, respectively. Besides nanoparticles shown above, there are also precipitates exhibiting a variety of morphologies and sizes as indicated in Fig. 11, electron diffraction experiment and EDS analysis prove that all the precipitates in S573 are Mg24Y5 phases. Moreover, the Mg3(Co, Y) phase transformed into small segments at this temperature,
3.2.4. Deformation behavior at 673 K The Mg matrix is significantly recrystallized along LPSO structures in S673, with dense SFs inside of the recrystallized grains, as shown in Fig. 12a and b. And the polygonal Mg24Y5 precipitates along the interface between the LPSO phase and Mg matrix are approximately 500 nm in diameter. Major Mg3(Co, Y) phase was broken into fragments, as shown in Fig. 12c. It should be mentioned that Mg3(Co, Y) particles were broken due to the dislocation shear in the S473 and S573, as shown in Fig. 6. In contrast, a large number of DRXed Mg grains in S673 might lead to fragmentation of the Mg3(Co, Y) particles [27]. A small number of deformation twins and kink bands were detected in the Mg matrix, as indicated in Fig. 12d. It should be mentioned that MgYCo4 phase still remains the same during deformation at this temperature due to its very large hardness. 4. Discussion 4.1. Deformation mechanism of Mg3(Co, Y) It's worthwhile to note that although the morphology shown in Fig. 6 is similar to the deformed γ-Y2Si2O7 ceramic [28] and Ti3Al intermetallic [29], but the deformation mechanisms are different. The high density of dislocations piles up in slip bands when the γ-Y2Si2O7 ceramic suffer great stress, inducing well-defined interlaced mesh amorphous layers [28]. And deformation twinning induced the grid architecture of the Ti3Al phase [29]. In contrast, the intersecting grid lines of the Mg3(Co, Y) phase exhibits dark contrast in Fig. 6a and very thin Mg3(Co, Y) crystalline exist in the grids in Fig. 6b, strongly indicating that the Mg3(Co, Y) phase coordinates the strain mainly 556
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Fig. 4. HAADF images of the {1012} twin and its interaction with intrinsic stacking fault (ISF) in S473 (the corresponding true strain is 22%). (a) {1012} deformation twins in Mg matrix; (b) A {1012} twin penetrating an ISF; (c) Corresponding HRSTEM image of panel (b). The twin boundary (TB) is marked by solid and dashed lines, which represent BP/PB boundary and coherent TB (CTB), respectively. (d) Periodic segregation of Co/Y in CTBs.
4.2. Precipitation of nanoscale Mg24Y5
through producing micro-cracks when deformed at 473 K. The study of the mechanism responsible for the micro-cracks along {1011} plane in Mg3(Co, Y) phase, rather than the {0001} plane, is ongoing now and will be published elsewhere.
Mg24Y5 precipitated in the samples during compression at 573 K, as shown in Fig. 8. Previous reports indicate that Mg24Y5 phase is the equilibrium phase of MgeY binary alloy. When MgeY alloys are heated at 573 K and above, Mg24Y5 phase can precipitate from the supersaturated solid solution [30] or transform into Mg24Y5 phase from the
Fig. 5. BF-TEM images showing deformation microstructure of the LPSO phases (the corresponding true strain is 22%). (a) Low magnification image indicating large amounts of dislocations were activated during the deformation process, as indicated by the arrow; (b) < a > dislocations or basal SFs gliding on basal planes, as indicated by the arrows; (c) Curves and kinks generated in a 15R-LPSO phase nearby a Mg3(Co, Y) phase, where “KB” represents “kink boundary”. 557
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Fig. 6. (a) Low and (b) high resolution HAADF-STEM images of the intermetallic Mg3(Co, Y) phase in S473 (the corresponding true strain is 22%), showing the micro-cracks propagate along {1101} planes of the Mg3(Co, Y) phase.
by the inserted image in Fig. 2a.
low-temperature phase [31]. We'll propose the formation mechanisms of Mg24Y5 phase in the following. Firstly, since the nano-sized Mg24Y5 phase in S573 distributed at the Mg matrix and micro-sized Mg24Y5 phase locates at the interface of LPSO/Mg phase or at the interface of LPSO/Mg3(Co, Y) phase at as-cast state, the possibility that the nanoscale Mg24Y5 phase formed during the casting process or transformed from the casting Mg24Y5 phase is ruled out here. Secondly, the nanoscale Mg24Y5 phase did not transform from the Mg3(Co, Y) phase. On one side, the Mg3(Co, Y) phase was just broken into small segments during the deformation. On the other side, if the nano-particles transform from the Mg3(Co, Y) phase, we are unable to explain where the Co element went because Co is almost insoluble in the Mg matrix. We suppose that the Mg24Y5 nano-particles precipitated during the deformation process. Our previous study showed that nanoparticles can precipitate in the matrix because supersaturated Y atoms segregate to the defects [32]. Further, we detected a segregation region of heavy elements at the LPSO/Mg interface at the as-cast sample, as indicated
4.3. Strengthening effect of intermetallic compounds in the Mg-Co-Y alloy From Table 2, we can conclude that the Mg-Co-Y alloy has high strength at all temperatures, especially at high temperatures. This mainly because of the strengthening effects caused by the various intermetallic compounds, such as LPSO, Mg24Y5, Mg3(Co, Y) and MgYCo4 phase, and their effects on the Mg matrix. Table 3 summarizes the corresponding deformation mechanisms of the Mg matrix and intermetallic phases in the Mg88Co5Y7 alloys during compression at different temperatures. At lower temperatures (373 K and 473 K), basal dislocations, deformation twins, and kink bands are responsible for the deformation of Mg matrix. Large amounts of basal slip nucleate in magnesium alloys since the critical shear stress (CRSS) of the basal slip is estimated to be about one-tenth of that of pyramidal slip [33]. {1012} tensile twin was activated in the Mg matrix of S473 since it is stable and
Fig. 7. Dislocations in Mg matrix of S573 (the corresponding true strain is 80%). (a) Dislocation walls parallel and vertical to basal plane of Mg matrix; (b) High resolution image of region ① in panel (a), showing a dissociated c + a dislocation; (c) High resolution image of region ② in panel (a), showing an array of a dislocations vertical to the basal planes; (d) FFT image of panel (c). BP represents “basal plane” and the observation direction is 〈1120〉. 558
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Fig. 8. Nanoparticles precipitated in the Mg matrix nearby LPSO phase in S573 (the corresponding true strain is 80%). (a) and (b) are BF-TEM image and the corresponding HAADF image, respectively.
compressed at 573 K. Therefore, the stress concentration between the phases can be released ingeniously although the hardness of each phase is different. No recrystallization in the Mg matrix at this temperature would effectively lower the recrystallization-induced softening effect for the alloy. The strengthening effect of Mg24Y5 precipitation might be negligible due to its distribution of not homogenous. In addition, MgYCo4 phases pin the matrix, while fragmented Mg3(Co, Y) phases and LPSO structures stabilize the grain growth, further offering the alloy high strength. The fragmentation of Mg3(Co, Y) phase may lead to the small strength decrease of the Mg88Co5Y7 alloy. As increasing of deformation temperature, the recrystallization was brought out and consequently results in softening in S673. However, the fragmentation of Mg3(Co, Y) phase and the grain boundary pinning effect of the hard LPSO and MgYCo4 phases make the compressive strengths of the Mg88Co5Y7 alloy at 573 K and 673 K higher than most of the magnesium alloys.
easy motivated, and kink bands were also detected. Deformation twins are generally considered to activate at low deformation temperature, while kink band form at relatively high temperature which doesn't favor the generation of twins (such as ≥573 K) [34,35]. Coexistence of deformation twins and kink bands in S473 is expected to relax more stress concentration and be beneficial for the ductility of the alloy. In S473, < a > dislocations and kink bands generated in some LPSO phase are helpful for the enhancement of strength and ductility of the Mg-Co-Y alloy. We did not detect recrystallized grains in the matrix of S573, although recrystallization always occurs in pure Mg and Mg alloys upon compressed at 573 K [36]. This can be illustrated from the following several aspects. First, the MgYCo4 phase within the matrix grain moves along with the deformation of Mg matrix, inducing little stress concentration on Mg/MgYCo4 interface. Second, the stress concentration on the interface of Mg matrix and LPSO phase can be released by precipitation of Mg24Y5 phase, where the Mg/LPSO interface accounts for the largest proportion of the whole alloy interface. Third, the stress concentration between Mg matrix and Mg3(Co, Y) phase can also be released when Mg3(Co, Y) phase transformed into small segments when
5. Conclusion The compressive properties of Mg88Co5Y7 alloy at 373–673 K are
Fig. 9. Morphology and corresponding EDS mapping images of the nanoparticles, indicating that the nanoparticles are enriched in Y elements, but no detectable Co elements. 559
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Fig. 10. (a–c) High-resolution HAADF-STEM images of the nanoparticles obtained from [001], [111] and [012] zone axis; (d–f) Corresponding FFT images of (a–c), respectively, demonstrating that the nanoparticles with an FCC structure should be Mg24Y5 phase combining the EDS results. The atomic models of Mg24Y5 phase observed from corresponding orientations are also exhibited in (a–c), where the orange circles represent Mg atoms and the green circles represent Y atoms. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 11. Other Mg24Y5 precipitates formed during 573 K deformation (the corresponding true strain is 80%).
only a small amount of Mg3(Co, Y) phase, while no deformation occurred in MgYCo4 intermetallic compound. 3. The LPSO phase, the MgYCo4 phase, and small Mg3(Co, Y) phases, to a certain extent, play an important role in pinning grain boundary and matrix and then enhance the strength of Mg-Co-Y alloy at higher temperatures (573 K and 673 K). A large number of dislocations containing c component were activated in the matrix, beneficial to the improvement of ductility for the sample deformed at 573 K. Recrystallization occurred in the most matrix grains compressed at 673 K, leading to a large strength reduction.
investigated and the microstructure evolution is characterized. The corresponding deformation mechanism of high strength and good plasticity of Mg88Co5Y7 alloy at various temperatures was discussed. The main conclusions are as follow: 1. The Mg88Co5Y7 alloy has good compressive mechanical properties at high temperatures. At 373 K and 473 K, the compressive strength is 311 MPa and the compression rate is 22%. The compressive strength at 573 K and 673 K are 234 MPa and 108 MPa, respectively, and both the compression rates are > 80%. 2. When the alloy deformed at lower temperatures (373 K and 473 K), a large number of {1012} tensile twins and deformation bands generated in the Mg matrix. Basal slip and deformation kink were motivated in the hard LPSO phase. Micro-cracks were formed in
Acknowledgments This work is supported by the National Natural Science Foundation 560
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Fig. 12. BF-TEM images of Mg88Co5Y7 alloy after compression at 673 K (the corresponding true strain is 80%). (a) Mg24Y5 phases precipitated at the interface of LPSO phase and Mg matrix; (b) Recrystallized fine grains with lots of SFs; (c) Mg3(Co, Y) phase turned into small segments and distributed along fringes of LPSO phase; (d) Deformation twins in un-recrystallized Mg matrix. Table 3 Deformation mechanisms of Mg matrix and intermetallic phases in Mg88Co5Y7 alloys during compression at varying temperatures. Temp.
373/473 K 573 K 673 K
Phases Mg matrix
LPSO
Mg3(Co, Y)
MgYCo4
a dislocation, twin c + a and a dislocations c + a and a dislocations, twin
a dislocation, kink, curve a dislocation, kink, curve a dislocation, kink, curve
Micro-crack Turned into segments Turned into segments
No change No change No change
of China (grants 51871222, 51801214, 51301177), the Innovation Fund of IMR (2015-PY08), and the Fund of SYNL (2015FP18, 2017FP16).
[10]
[11]
References [1] M. Yamasaki, T. Anan, S. Yoshimoto, Y. Kawamura, Mechanical properties of warmextruded Mg–Zn–Gd alloy with coherent 14H long periodic stacking ordered structure precipitate, Scr. Mater. 53 (7) (2005) 799–803. [2] Y. Kawamura, T. Kasahara, S. Izumi, M. Yamasaki, Elevated temperature Mg97Y2Cu1 alloy with long period ordered structure, Scr. Mater. 55 (5) (2006) 453–456. [3] Y. Kawamura, M. Yamasaki, Formation and mechanical properties of Mg97Zn1RE2 alloys with long-period stacking ordered structure, Mater. Trans. 48 (11) (2007) 2986–2992. [4] T. Itoi, K. Takahashi, H. Moriyama, M. Hirohashi, A high-strength Mg–Ni–Y alloy sheet with a long-period ordered phase prepared by hot-rolling, Scr. Mater. 59 (10) (2008) 1155–1158. [5] S. Yoshimoto, M. Yamasaki, Y. Kawamura, Microstructure and mechanical properties of extruded Mg-Zn-Y alloys with 14H long period ordered structure, Mater. Trans. 47 (4) (2006) 959–965. [6] T. Itoi, T. Inazawa, M. Yamasaki, Y. Kawamura, M. Hirohashi, Microstructure and mechanical properties of Mg–Zn–Y alloy sheet prepared by hot-rolling, Mater. Sci. Eng. A 560 (2013) 216–223. [7] X.H. Shao, Z.Q. Yang, X.L. Ma, Strengthening and toughening mechanisms in Mg–Zn–Y alloy with a long period stacking ordered structure, Acta Mater. 58 (14) (2010) 4760–4771. [8] Z.Z. Peng, X.H. Shao, Q.Q. Jin, J.F. Liu, X.L. Ma, Dislocation configuration and solute redistribution of low angle kink boundaries in an extruded Mg–Zn–Y–Zr alloy, Mater. Sci. Eng. A 687 (2017) 211–220. [9] M. Matsuda, S. Ii, Y. Kawamura, Y. Ikuhara, M. Nishida, Variation of long-period
[12]
[13] [14] [15] [16]
[17] [18] [19]
[20] [21]
561
Precipitation No Mg24Y5 Mg24Y5
stacking order structures in rapidly solidified Mg97Zn1Y2 alloy, Mater. Sci. Eng. A 393 (1–2) (2005) 269–274. Y.M. Zhu, M. Weyland, A.J. Morton, K. Oh-ishi, K. Hono, J.F. Nie, The building block of long-period structures in Mg–RE–Zn alloys, Scr. Mater. 60 (11) (2009) 980–983. X.H. Shao, Z.Q. Yang, X.L. Ma, Interplay between deformation twins and basal stacking faults enriched with Zn/Y in Mg97Zn1Y2 alloy, Philos. Mag. Lett. 94 (3) (2014) 150–156. X.H. Shao, S.J. Zheng, D. Chen, Q.Q. Jin, Z.Z. Peng, X.L. Ma, Deformation twinning induced decomposition of lamellar LPSO structure and its re-precipitation in an MgZn-Y alloy, Sci. Rep. 6 (2016) 30096. W.W. Hu, Z.Q. Yang, H.Q. Ye, < c+a > dislocations and their interactions with other crystal defects in a Mg alloy, Acta Mater. 124 (2017) 372–382. S.-B. Mi, Q.-Q. Jin, New polytypes of long-period stacking ordered structures in Mg–Co–Y alloys, Scr. Mater. 68 (8) (2013) 635–638. Q.Q. Jin, X.H. Shao, X.B. Hu, Z.Z. Peng, X.L. Ma, New polytypes of LPSO structures in an Mg–Co–Y alloy, Philos. Mag. 97 (1) (2016) 1–16. Q.Q. Jin, X.H. Shao, Z.Z. Peng, J.H. You, K.Q. Qiu, X.L. Ma, Crystallographic account of an ultra-long period stacking ordered phase in an Mg88Co5Y7 alloy, J. Alloys Compd. 693 (2017) 1035–1038. Q.-Q. Jin, S.-B. Mi, Intermetallic phases in Mg–Co–Y alloys, J. Alloys Compd. 582 (2014) 130–134. M. Egami, E. Abe, Structure of a novel Mg-rich complex compound in Mg–Co–Y ternary alloys, Scr. Mater. 98 (2015) 64–67. Z. Yang, M.F. Chisholm, G. Duscher, X. Ma, S.J. Pennycook, Direct observation of dislocation dissociation and Suzuki segregation in a Mg–Zn–Y alloy by aberrationcorrected scanning transmission electron microscopy, Acta Mater. 61 (1) (2013) 350–359. X. Shao, Characterization of Long Period Stacking Ordered Phases and Their Effects on Mechanical Properties of Mg Alloys, Insititute of Metal Research Doctor, 2011. B.-J. Lv, J. Peng, Y. Peng, A.-T. Tang, F.-S. Pan, The effect of LPSO phase on hot deformation behavior and dynamic recrystallization evolution of Mg–2.0Zn–0.3Zr–5.8Y alloy, Mater. Sci. Eng. A 579 (2013) 209–216.
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Z.Z. Peng, et al.
[29] K. Kishida, Y. Takahama, H. Inui, c-axis compression twinning in an off-stoichiometric compound Ti3Al with the D019 structure, Mater. Sci. Eng. A 400-401 (2005) 339–344. [30] M.-X. Zhang, P.M. Kelly, Morphology and crystallography of Mg24Y5 precipitate in Mg–Y alloy, Scr. Mater. 48 (2003) 379–384. [31] D. Mizer, B.C. Perters, A study of precipitation at elevated temperatures in a Mg-8.7 pct Y alloy, Metall. Trans. 3 (1972) 3262–3264. [32] Z.Z. Peng, X.H. Shao, Q.Q. Jin, C.H. Li, X.L. Ma, Deformation induced precipitation of nano-particles in an Mg88Co5Y7 alloy, Scr. Mater. 116 (2016) 57–61. [33] N. Stanford, R. Cottam, B. Davis, J. Robson, Evaluating the effect of yttrium as a solute strengthener in magnesium using in situ neutron diffraction, Acta Mater. 78 (2014) 1–13. [34] X. Yang, Y. Jiang, Morphology and crystallographic characteristics of deformation bands in Mg alloy under hot deformation, Acta Metall. Sin. 46 (4) (2010) 451–457. [35] L.X. Wang, G. Fang, M.A. Leeflang, J. Duszczyk, J. Zhou, Constitutive behavior and microstructure evolution of the as-extruded AE21 magnesium alloy during hot compression testing, J. Alloys Compd. 622 (2015) 121–129. [36] B.-J. Lv, J. Peng, Y.-J. Wang, X.-Q. An, L.-P. Zhong, A.-T. Tang, F.-S. Pan, Dynamic recrystallization behavior and hot workability of Mg–2.0Zn–0.3Zr–0.9Y alloy by using hot compression test, Mater. Des. 53 (2014) 357–365.
[22] B.-J. Lv, J. Peng, L.-L. Zhu, Y.-J. Wang, A.-T. Tang, The effect of 14H LPSO phase on dynamic recrystallization behavior and hot workability of Mg–2.0Zn–0.3Zr–5.8Y alloy, Mater. Sci. Eng. A 599 (2014) 150–159. [23] X. Xia, Q. Chen, J. Li, D. Shu, C. Hu, S. Huang, Z. Zhao, Characterization of hot deformation behavior of as-extruded Mg–Gd–Y–Zn–Zr alloy, J. Alloys Compd. 610 (2014) 203–211. [24] N. Tahreen, D.F. Zhang, F.S. Pan, X.Q. Jiang, C. Li, D.Y. Li, D.L. Chen, Characterization of hot deformation behavior of an extruded Mg–Zn–Mn–Y alloy containing LPSO phase, J. Alloys Compd. 644 (2015) 814–823. [25] J.W. Christian, S. Mahajan, Deformation twinning, Prog. Mater. Sci. 39 (1–2) (1995) 1–157. [26] K. Máthis, K. Nyilas, A. Axt, I. Dragomir-Cernatescu, T. Ungár, P. Lukáč, The evolution of non-basal dislocations as a function of deformation temperature in pure magnesium determined by X-ray diffraction, Acta Mater. 52 (10) (2004) 2889–2894. [27] P. Perez, M. Eddahbi, S. Gonzalez, G. Garces, P. Adeva, Refinement of the microstructure during superplastic deformation of extruded Mg94Ni3Y1.5CeMM1.5 alloy, Scr. Mater. 64 (1) (2011) 33–36. [28] Z.J. Lin, M.J. Zhuo, Z.Q. Sun, P. Veyssiere, Y.C. Zhou, Amorphization by dislocation accumulation in shear bands, Acta Mater. 57 (9) (2009) 2851–2857.
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Effect of temperature on deformation mechanisms of the Mg88Co5Y7 alloy during hot compression ⁎
Z.Z. Penga, Q.Q. Jina, X.H. Shaoa, , Y.T. Zhoua, S.J. Zhenga, B. Zhanga, X.L. Maa,b, a b
T
⁎
Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, 110016 Shenyang, China School of Materials Science and Engineering, Lanzhou University of Technology, 287 Langongping Road, 730050 Lanzhou, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Magnesium alloy LPSO structure Deformation twin Recrystallization Intermetallics
The deformation mechanisms and microstructure evolution of Mg88Co5Y7 (at.%) alloy with long period stacking ordered (LPSO) phase during compression at 373–673 K have been studied using transmission electron microscopy. The alloy consists of α-Mg matrix, interdendritic LPSO phase, and other Mg-Co-Y intermetallic compounds. Increasing temperature leads to a general decrease in strength of the Mg88Co5Y7 alloy. A large number of {1012 } tension twins and deformation bands are activated in the matrix, whereas basal slip and deformation kink dominate in the LPSO structures during compression at 373 and 473 K. Non-basal slip in matrix and almost no dynamic recrystallization are responsible for high strength and the good ductility of sample deformed at 573 K. Recrystallization of Mg matrix occurs upon deformation at 673 K, dramatically lowering the corresponding strength. Further, the grain boundary pinning effect from LPSO phase, MgYCo4 phase, and broken Mg3(Co, Y) segments is supposed to account for the relatively high strength of Mg88Co5Y7 alloy at high temperatures.
1. Introduction Because of low strength and poor plasticity, commercial magnesium alloys as structural materials have limited application in the automobile industry and other industries. Mg-TM-RE (TM, transition metals; RE, rare earth elements) ternary alloys containing long period stacking ordered (LPSO) phases have received considerable attention during the past decades due to their excellent performance both at room temperature and elevated temperatures [1–8]. Of the ternary alloy systems that LPSO phases have been observed, Mg-TM-RE (TM = Ni, Cu and Zn; RE = Y, Gd, Tb, Dy, Ho, and Er) alloys [2–4,9] were reported to contain LPSO phases with AB′C′A stacking faults, where B′ and C′ layers are commonly enriched by TM/RE atoms [10]. The mechanical properties of these alloys consisting of Mg and LPSO phases have been extensively studied [1–7]. The deformation mechanisms include deformation kink [7], deformation twin and its interaction with LPSO phases [11,12] and non-basal dislocations in matrix [8,13]. Very recently, new polytypes of LPSO structures featuring AB′C stacking faults (B′ layer is enriched in Co/Y atoms), as well as some other intermetallic compounds, such as Mg24Y5, Mg3(Co, Y) and MgYCo4 phases, were characterized in Mg-Co-Y as-cast alloys [14–16]. The microstructural details (including morphology, crystal structure and volume fraction of the constituent phase) and phase transformation of the as-cast and annealed Mg88Co5Y7 (at.%) alloys ⁎
have been elaborated [17,18]. However, although the hardness measurements demonstrate that the LPSO structures in Mg-Co-Y systems are harder than those in Mg-Zn-Y alloys [7,14], there is very little information considering mechanical properties and related deformation mechanisms of Mg-Co-Y alloys containing AB′C-LPSO structure and other intermetallic compounds. In the present paper, we investigated the deformation behaviors of the as-cast Mg88Co5Y7 (at.%) alloy during hot compression. The microstructure evolution of the samples at different temperatures was characterized by scanning electron microscope (SEM) and transmission electron microscope (TEM). The corresponding mechanisms are also discussed. Further, we try to reveal the effects of the LPSO phase and intermetallic on the strength and ductility of the Mg88Co5Y7 alloy, which will shed some new light on optimizing the mechanical properties of the magnesium alloys by introducing distinctive microstructures. 2. Experimental procedures A ternary alloy with a nominal composition Mg88Co5Y7 (at.%) was prepared by melting the high purity Mg, Co and Mg-30wt%Y master alloys in a graphite crucible under protection of ultrahigh purity argon atmosphere in a high-frequency induction melting furnace, and then cooled down to the room temperature under the argon atmosphere in the furnace. Specimens with the dimensions 4 × 4 × 8 mm3 were cut
Corresponding authors. E-mail addresses: [email protected] (X.H. Shao), [email protected] (X.L. Ma).
https://doi.org/10.1016/j.matchar.2019.03.049 Received 10 December 2018; Received in revised form 9 March 2019; Accepted 31 March 2019 Available online 01 April 2019 1044-5803/ © 2019 Elsevier Inc. All rights reserved.
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from the ingot by low-speed diamond cutting machine cooling by water. Compression experiment was carried out at 373 K, 473 K, 573 K, and 673 K and a strain rate of 1.0 × 10−3 s−1 in a Gleeble-3800 thermal simulation machine. We refer these samples S373, S473, S573, and S673 hereafter, respectively. Prior to being compressed the specimens were conductively heated to the setting temperatures at a heating rate of 5 K s−1 and held for 180 s for equilibration. The compression direction was parallel to the long axis of the specimens. The specimens were deformed to failure (S373 and S473) or stopped at the setting strain (80% for S573 and S673), then immediately water quenched to room temperature for microstructure analysis. The microstructures of specimens were examined by scanning electron microscopy (SEM) and scanning transmission electron microscopy (STEM). Thin foil samples for TEM were prepared by the conventional ion milling method, a Gatan precision ion polishing system (PIPS 691) with a liquid-nitrogen-cooled stage to avoid preferential thinning effects. The selected area electron diffraction (SAED) patterns and bright-field (BF) TEM images were obtained from the JEOL JEM 2100F FEG-TEM operating at 200 kV. The Titan3™ G2 60-300, which was operated at 300 kV and equipped with an energy-dispersive X-ray (EDX) detector and a high-angle annular dark-field (HAADF) detector, was used to obtain high-resolution STEM images and to measure chemical compositions.
Table 1 Compressive stress, yield stress and a compression ratio of Mg88Co5Y7 alloy at 373 K, 473 K, 573 K, and 673 K respectively. Compression stress σ (MPa)
Yield stress σ0.2 (MPa)
Compression ratio δ (%)
311 311 234 108
145 145 100 90
22 22 > 80 > 80
373 K 473 K 573 K 673 K
Table 2 Compressive stresses of Mg88Co5Y7 (at.%) alloys and other LPSO containing magnesium alloys. (The compressive stresses here are true stresses from the true stress-strain curves and the strain rate are all 1.0 × 10−3 s−1.) Magnesium alloys
Mg97Zn1Y2 (at.%) MgZn2Zr0.3Y5.8 (wt%) 693 K, 12 h MgZn2Zr0.3Y5.8 (wt%) 773 K, 20 h Mg-Gd-Y-Zn-Zr Mg-Zn-Mn-Y Mg88Co5Y7 (at.%)
Compression stress σ (MPa)
Ref.
373 K
473 K
573 K
673 K
223
217
311
317 311
190 164 210 213 246 234
96 75 89 50 79 108
[19] [20] [21] [22] [23] Present
This highlights that the Mg88Co5Y7 (at.%) alloy shows high compressive stress at high temperatures.
3. Results 3.1. True stress–true strain curves
other magnesium alloys listed in Table 2, we can conclude that the compressive strength of Mg88Co5Y7 alloy is superior to most cast [20] and homogenized LPSO containing magnesium alloys [21,22], nearly equivalent to the mechanical strength of extruded LPSO containing magnesium alloys [23,24], especially at 573 K and 673 K.
The compression stress–strain curves of the Mg88Co5Y7 alloy under a strain rate of 1.0 × 10−3 s−1 at 373 K, 473 K, 573 K, and 673 K are shown in Fig. 1. The curves for S373 and S473 are nearly coincident, with the peak stress of about 311 MPa and the compression ratio of 22%, suggesting that the deformation behavior at 373 K and 473 K should be similar. The strength for S573 decreases by 24.8% (from 311 MPa to 234 MPa), whereas the compression ratio improves over 264% (from 22% to above 80%). In contrast, the strength of S673 dramatically decreases to 108 MPa. The serrated flow may suggest dynamic strain aging due to the mutual interaction between dislocations and solute atoms during hot compression since solute atoms in Mg alloys containing LPSO structures tend to segregate along dislocations and stacking faults [19]. But the underlying mechanism would be investigated in the near future. Table 1 summarizes the compressive strength, yield strength (σ0.2) and a compression ratio of the Mg88Co5Y7 samples compressed at different temperatures. Compared with the compression properties of
3.2. Microstructural evolution and deformation mechanisms of Mg88Co5Y7 deformed alloy 3.2.1. Microstructural morphologies of Mg88Co5Y7 alloy deformed at different temperatures Fig. 2 are the back-scattering electron (BSE) SEM image of the Mg88Co5Y7 alloy, clearly presenting the evolution of intermetallic compounds during compression. As indicated by the small arrows in Fig. 2a, four kinds of intermetallics, LPSO phase, Mg24Y5 phase (about 0.9% volume fraction), Mg3(Co, Y) phase and MgYCo4 phase, coexist at the as-cast state. The details about their composition, crystal structure and volume fraction in the cast alloy have been illustrated in refs. [14, 17, 18]. Fig. 2b–d presents the deformation microstructures of S473, S573, and S673, respectively. Fig. 2b shows that the morphology of S473 is almost the same as that at as-cast state, except some of the LPSO phases are slightly curved. The S373 is structurally analogous to S473, which is not shown. For S573, Fig. 2c indicates that the intergrown LPSO/Mg3(Co, Y) phase is obviously curved, which may infer the formation of deformation kink. Fig. 2d shows that Mg matrix with LPSO phase and Mg3(Co, Y) phase at the grain boundary were elongated perpendicular to the loading axis during 673 K. Moreover, the Mg3(Co, Y) phase evolved into small particles and dispersed along the edges of the LPSO phases, as marked by the yellow arrows in the inserted zoomin image in Fig. 2d. MgYCo4 phase in Mg matrix seems to be undeformed during compression. In order to further illustrate the corresponding deformation mechanisms of Mg88Co5Y7 alloy at various temperatures, we studied the samples by TEM observation. Note that in this work the observation plane of the TEM sample is parallel to the compression axis.
Fig. 1. True stress–true strain curves obtained from compressive tests at 373 K, 473 K, 573 K, and 673 K under a constant strain rate of 1.0 × 10−3 s−1 respectively.
3.2.2. Deformation behavior at 473 K Profuse deformation twins and kinks occur in Mg matrix of S473. Each of the five bands in Fig. 3a, numbered “1–5”, exhibits twin 554
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Fig. 2. Back-scattering electron (BSE) images of the Mg88Co5Y7 alloys. (a) As-cast (The arrows in the enlarged area show the segregation regions of heavy elements at the interface between LPSO phase and Mg matrix); (b) S473; (c) S573 and (d) S673, showing the evolution of the microstructures. (For interpretation of the references to color in this figure, the reader is referred to the web version of this article.)
between {1012} twin and ISF here is consistent with that in compressed Mg97Zn1Y2 alloys [11]. Meanwhile, Co and/or Y atoms periodically segregated along the CTB a little bit farther from the intersection between twin and ISF, as shown in Fig. 4d. Additionally, Fig. 4 also demonstrated that heavy solute atoms are easily segregated to stacking faults and twin boundaries in this alloy. Basal dislocations, low and high angle grain boundaries (GBs) were produced in the LPSO phases. Fig. 5a is the BF-TEM image of one deformed LPSO phase, indicating large amounts of dislocations are activated during the deformation process. The dark line contrast along basal plane in the BF-TEM image acquired along the 〈1120〉 orientation of the LPSO phase (Fig. 5b), suggesting that < a > dislocations or basal SFs glide on basal planes, as indicated by the arrows. The accumulation of the defects described above can make the LPSO slightly bent and even result in deformation kink [8]. For example, multiple kinks in the LPSO phase were detected when the local stress increased at the interface between the 15R-LPSO and a harder intermetallic Mg3(Co, Y) phase, as shown in Fig. 5c. A very small amount of Mg3(Co, Y) phase was broken. Fig. 6a is a low magnification HAADF image of the intermetallic Mg3(Co, Y) in S473, showing the cracks propagate along the {1101} planes of the Mg3(Co, Y) phase. The Co2Y9 clusters in Mg3 (Co, Y) cell in the same row deflected from each other along the {1011} plane, as shown in Fig. 6b, where the fast Fourier transformation (FFT) image is inserted. Combining the crystal structure, we assume that the fracture plane of Mg3(Co, Y) is closely related to the chemical bonding of the atoms, which need further related simulation.
orientation with its neighboring band. The SAED pattern of the circle area in Fig. 3a indicates that the “1” and its neighbors show {1012} tensile twin orientation relationship, demonstrated by electron diffraction (ED) pattern in Fig. 3b. Actually, {1012} tensile twin is the main twin in the deformed Mg matrix under this deformation condition in our study. Fig. 3c and d is the BF-TEM and HAADF images of the deformation bands in the Mg matrix. They always show sharp interfaces with the Mg matrix, similar to the morphology of twins. Nevertheless, the deformation band always deviates from the zone axis when the matrix is rotated along 〈1120〉 or 〈1010〉 direction. This can be further proved by the strong contrast differences between the deformation band and the matrix in the BF-TEM and HAADF images. These deformation bands might be other twin bands rather than traditional {1012}, {1011} or {1013} twin bands [25], which will be explored in the near future. It should be mentioned the multiple twins and deformation bands are always generated together, crossing or paralleling to each other, as shown in Fig. 3a and c. The migrating of the twin boundary (TB) and its interaction with other defects are also frequently detected in S473. Fig. 4a and b shows HAADF images of the {1012} twin and its interaction with the intrinsic stacking fault (ISF), respectively. Fig. 4c is the corresponding high-resolution STEM image of the framed area in Fig. 4b. The TB consists of BP/PB boundary (B represents basal plane, P represents prismatic plane) and coherent twin boundary (CTB), respectively represented by solid and dashed lines. The stacking sequence of the ISF is “…ABAB′CBC…”, where the B′ layer is brighter due to a higher concentration of Co and/or Y elements compared with the adjacent magnesium layers. The stacking structure of AB′C is same as the stacking block of 15R, 12H and 21R-LPSO in the Mg-Co-Y alloy [14]. After twinning, the original AB′C fault structure converts to the ABAB stacking sequence in the twin, as shown in Fig. 4c. The interaction
3.2.3. Deformation behavior at 573 K Profuse dislocations in the Mg matrix were detected in S573, as shown in Fig. 7a. The pyramidal dislocation movement has more 555
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Fig. 3. Twins and deformation bands formed in the deformed Mg matrix. (a) and (b) Morphology of the twins and selected area electron diffraction (SAED) pattern of the {1012} twin; (c) and (d) BF-TEM and HAADF images of the deformation band, respectively.
as shown in Fig. 8.
advantages than twinning motivation in energy when the deformation temperature is above 200 °C [26]. Fig. 7b and c is the high-resolution TEM images of the formed boundary at location ① and ②, respectively, suggesting that the boundary parallel or vertical to the Mg basal plane is composed of the dissociated < c + a > dislocations and an array of < a > dislocations. Fig. 7d is the FFT pattern of Fig. 7c, indicating the small rotation angle of the basal planes is 3.6°. In contrast, deformation twins or kink bands were rarely observed in the matrix. The most prominent feature in S573 is the precipitation of nanoparticles in the Mg matrix nearby the LPSO phase, as shown in Fig. 8. Fig. 8a and b shows BF-TEM and corresponding HAADF image of the nanoparticles, respectively, indicating that the particles with polyhedron shape are in the range of 100–300 nm. To clarify the nanoparticles, the chemical composition and crystal structure of the particle were analyzed. Fig. 9 shows morphology and corresponding EDS mapping images of the nanoparticles, indicating that the nanoparticles are enriched in Y element and no detectable Co element. The specific chemical composition of the nanoparticle is estimated to be 88.99 ± 0.95 at.% Mg – 0.04 ± 0.02 at.% Co – 10.96 ± 0.97 at.% Y, which suggests that it should be MgeY binary phase. Fig. 10a–c is highresolution STEM images of the nanoparticles observing from various zone axis, and the corresponding FFT images of Fig. 10a–c are shown in Fig. 10d–f, respectively. Combined the comprehensive analysis of chemical composition and crystal structure, it's fair to say that the nanoparticle is Mg24Y5 phase. The atomic models of Mg24Y5 phase are superimposed in Fig. 10a–c, where blue and orange dots represent Y and Mg atoms, respectively. Besides nanoparticles shown above, there are also precipitates exhibiting a variety of morphologies and sizes as indicated in Fig. 11, electron diffraction experiment and EDS analysis prove that all the precipitates in S573 are Mg24Y5 phases. Moreover, the Mg3(Co, Y) phase transformed into small segments at this temperature,
3.2.4. Deformation behavior at 673 K The Mg matrix is significantly recrystallized along LPSO structures in S673, with dense SFs inside of the recrystallized grains, as shown in Fig. 12a and b. And the polygonal Mg24Y5 precipitates along the interface between the LPSO phase and Mg matrix are approximately 500 nm in diameter. Major Mg3(Co, Y) phase was broken into fragments, as shown in Fig. 12c. It should be mentioned that Mg3(Co, Y) particles were broken due to the dislocation shear in the S473 and S573, as shown in Fig. 6. In contrast, a large number of DRXed Mg grains in S673 might lead to fragmentation of the Mg3(Co, Y) particles [27]. A small number of deformation twins and kink bands were detected in the Mg matrix, as indicated in Fig. 12d. It should be mentioned that MgYCo4 phase still remains the same during deformation at this temperature due to its very large hardness. 4. Discussion 4.1. Deformation mechanism of Mg3(Co, Y) It's worthwhile to note that although the morphology shown in Fig. 6 is similar to the deformed γ-Y2Si2O7 ceramic [28] and Ti3Al intermetallic [29], but the deformation mechanisms are different. The high density of dislocations piles up in slip bands when the γ-Y2Si2O7 ceramic suffer great stress, inducing well-defined interlaced mesh amorphous layers [28]. And deformation twinning induced the grid architecture of the Ti3Al phase [29]. In contrast, the intersecting grid lines of the Mg3(Co, Y) phase exhibits dark contrast in Fig. 6a and very thin Mg3(Co, Y) crystalline exist in the grids in Fig. 6b, strongly indicating that the Mg3(Co, Y) phase coordinates the strain mainly 556
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Fig. 4. HAADF images of the {1012} twin and its interaction with intrinsic stacking fault (ISF) in S473 (the corresponding true strain is 22%). (a) {1012} deformation twins in Mg matrix; (b) A {1012} twin penetrating an ISF; (c) Corresponding HRSTEM image of panel (b). The twin boundary (TB) is marked by solid and dashed lines, which represent BP/PB boundary and coherent TB (CTB), respectively. (d) Periodic segregation of Co/Y in CTBs.
4.2. Precipitation of nanoscale Mg24Y5
through producing micro-cracks when deformed at 473 K. The study of the mechanism responsible for the micro-cracks along {1011} plane in Mg3(Co, Y) phase, rather than the {0001} plane, is ongoing now and will be published elsewhere.
Mg24Y5 precipitated in the samples during compression at 573 K, as shown in Fig. 8. Previous reports indicate that Mg24Y5 phase is the equilibrium phase of MgeY binary alloy. When MgeY alloys are heated at 573 K and above, Mg24Y5 phase can precipitate from the supersaturated solid solution [30] or transform into Mg24Y5 phase from the
Fig. 5. BF-TEM images showing deformation microstructure of the LPSO phases (the corresponding true strain is 22%). (a) Low magnification image indicating large amounts of dislocations were activated during the deformation process, as indicated by the arrow; (b) < a > dislocations or basal SFs gliding on basal planes, as indicated by the arrows; (c) Curves and kinks generated in a 15R-LPSO phase nearby a Mg3(Co, Y) phase, where “KB” represents “kink boundary”. 557
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Fig. 6. (a) Low and (b) high resolution HAADF-STEM images of the intermetallic Mg3(Co, Y) phase in S473 (the corresponding true strain is 22%), showing the micro-cracks propagate along {1101} planes of the Mg3(Co, Y) phase.
by the inserted image in Fig. 2a.
low-temperature phase [31]. We'll propose the formation mechanisms of Mg24Y5 phase in the following. Firstly, since the nano-sized Mg24Y5 phase in S573 distributed at the Mg matrix and micro-sized Mg24Y5 phase locates at the interface of LPSO/Mg phase or at the interface of LPSO/Mg3(Co, Y) phase at as-cast state, the possibility that the nanoscale Mg24Y5 phase formed during the casting process or transformed from the casting Mg24Y5 phase is ruled out here. Secondly, the nanoscale Mg24Y5 phase did not transform from the Mg3(Co, Y) phase. On one side, the Mg3(Co, Y) phase was just broken into small segments during the deformation. On the other side, if the nano-particles transform from the Mg3(Co, Y) phase, we are unable to explain where the Co element went because Co is almost insoluble in the Mg matrix. We suppose that the Mg24Y5 nano-particles precipitated during the deformation process. Our previous study showed that nanoparticles can precipitate in the matrix because supersaturated Y atoms segregate to the defects [32]. Further, we detected a segregation region of heavy elements at the LPSO/Mg interface at the as-cast sample, as indicated
4.3. Strengthening effect of intermetallic compounds in the Mg-Co-Y alloy From Table 2, we can conclude that the Mg-Co-Y alloy has high strength at all temperatures, especially at high temperatures. This mainly because of the strengthening effects caused by the various intermetallic compounds, such as LPSO, Mg24Y5, Mg3(Co, Y) and MgYCo4 phase, and their effects on the Mg matrix. Table 3 summarizes the corresponding deformation mechanisms of the Mg matrix and intermetallic phases in the Mg88Co5Y7 alloys during compression at different temperatures. At lower temperatures (373 K and 473 K), basal dislocations, deformation twins, and kink bands are responsible for the deformation of Mg matrix. Large amounts of basal slip nucleate in magnesium alloys since the critical shear stress (CRSS) of the basal slip is estimated to be about one-tenth of that of pyramidal slip [33]. {1012} tensile twin was activated in the Mg matrix of S473 since it is stable and
Fig. 7. Dislocations in Mg matrix of S573 (the corresponding true strain is 80%). (a) Dislocation walls parallel and vertical to basal plane of Mg matrix; (b) High resolution image of region ① in panel (a), showing a dissociated c + a dislocation; (c) High resolution image of region ② in panel (a), showing an array of a dislocations vertical to the basal planes; (d) FFT image of panel (c). BP represents “basal plane” and the observation direction is 〈1120〉. 558
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Fig. 8. Nanoparticles precipitated in the Mg matrix nearby LPSO phase in S573 (the corresponding true strain is 80%). (a) and (b) are BF-TEM image and the corresponding HAADF image, respectively.
compressed at 573 K. Therefore, the stress concentration between the phases can be released ingeniously although the hardness of each phase is different. No recrystallization in the Mg matrix at this temperature would effectively lower the recrystallization-induced softening effect for the alloy. The strengthening effect of Mg24Y5 precipitation might be negligible due to its distribution of not homogenous. In addition, MgYCo4 phases pin the matrix, while fragmented Mg3(Co, Y) phases and LPSO structures stabilize the grain growth, further offering the alloy high strength. The fragmentation of Mg3(Co, Y) phase may lead to the small strength decrease of the Mg88Co5Y7 alloy. As increasing of deformation temperature, the recrystallization was brought out and consequently results in softening in S673. However, the fragmentation of Mg3(Co, Y) phase and the grain boundary pinning effect of the hard LPSO and MgYCo4 phases make the compressive strengths of the Mg88Co5Y7 alloy at 573 K and 673 K higher than most of the magnesium alloys.
easy motivated, and kink bands were also detected. Deformation twins are generally considered to activate at low deformation temperature, while kink band form at relatively high temperature which doesn't favor the generation of twins (such as ≥573 K) [34,35]. Coexistence of deformation twins and kink bands in S473 is expected to relax more stress concentration and be beneficial for the ductility of the alloy. In S473, < a > dislocations and kink bands generated in some LPSO phase are helpful for the enhancement of strength and ductility of the Mg-Co-Y alloy. We did not detect recrystallized grains in the matrix of S573, although recrystallization always occurs in pure Mg and Mg alloys upon compressed at 573 K [36]. This can be illustrated from the following several aspects. First, the MgYCo4 phase within the matrix grain moves along with the deformation of Mg matrix, inducing little stress concentration on Mg/MgYCo4 interface. Second, the stress concentration on the interface of Mg matrix and LPSO phase can be released by precipitation of Mg24Y5 phase, where the Mg/LPSO interface accounts for the largest proportion of the whole alloy interface. Third, the stress concentration between Mg matrix and Mg3(Co, Y) phase can also be released when Mg3(Co, Y) phase transformed into small segments when
5. Conclusion The compressive properties of Mg88Co5Y7 alloy at 373–673 K are
Fig. 9. Morphology and corresponding EDS mapping images of the nanoparticles, indicating that the nanoparticles are enriched in Y elements, but no detectable Co elements. 559
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Fig. 10. (a–c) High-resolution HAADF-STEM images of the nanoparticles obtained from [001], [111] and [012] zone axis; (d–f) Corresponding FFT images of (a–c), respectively, demonstrating that the nanoparticles with an FCC structure should be Mg24Y5 phase combining the EDS results. The atomic models of Mg24Y5 phase observed from corresponding orientations are also exhibited in (a–c), where the orange circles represent Mg atoms and the green circles represent Y atoms. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 11. Other Mg24Y5 precipitates formed during 573 K deformation (the corresponding true strain is 80%).
only a small amount of Mg3(Co, Y) phase, while no deformation occurred in MgYCo4 intermetallic compound. 3. The LPSO phase, the MgYCo4 phase, and small Mg3(Co, Y) phases, to a certain extent, play an important role in pinning grain boundary and matrix and then enhance the strength of Mg-Co-Y alloy at higher temperatures (573 K and 673 K). A large number of dislocations containing c component were activated in the matrix, beneficial to the improvement of ductility for the sample deformed at 573 K. Recrystallization occurred in the most matrix grains compressed at 673 K, leading to a large strength reduction.
investigated and the microstructure evolution is characterized. The corresponding deformation mechanism of high strength and good plasticity of Mg88Co5Y7 alloy at various temperatures was discussed. The main conclusions are as follow: 1. The Mg88Co5Y7 alloy has good compressive mechanical properties at high temperatures. At 373 K and 473 K, the compressive strength is 311 MPa and the compression rate is 22%. The compressive strength at 573 K and 673 K are 234 MPa and 108 MPa, respectively, and both the compression rates are > 80%. 2. When the alloy deformed at lower temperatures (373 K and 473 K), a large number of {1012} tensile twins and deformation bands generated in the Mg matrix. Basal slip and deformation kink were motivated in the hard LPSO phase. Micro-cracks were formed in
Acknowledgments This work is supported by the National Natural Science Foundation 560
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Fig. 12. BF-TEM images of Mg88Co5Y7 alloy after compression at 673 K (the corresponding true strain is 80%). (a) Mg24Y5 phases precipitated at the interface of LPSO phase and Mg matrix; (b) Recrystallized fine grains with lots of SFs; (c) Mg3(Co, Y) phase turned into small segments and distributed along fringes of LPSO phase; (d) Deformation twins in un-recrystallized Mg matrix. Table 3 Deformation mechanisms of Mg matrix and intermetallic phases in Mg88Co5Y7 alloys during compression at varying temperatures. Temp.
373/473 K 573 K 673 K
Phases Mg matrix
LPSO
Mg3(Co, Y)
MgYCo4
a dislocation, twin c + a and a dislocations c + a and a dislocations, twin
a dislocation, kink, curve a dislocation, kink, curve a dislocation, kink, curve
Micro-crack Turned into segments Turned into segments
No change No change No change
of China (grants 51871222, 51801214, 51301177), the Innovation Fund of IMR (2015-PY08), and the Fund of SYNL (2015FP18, 2017FP16).
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References [1] M. Yamasaki, T. Anan, S. Yoshimoto, Y. Kawamura, Mechanical properties of warmextruded Mg–Zn–Gd alloy with coherent 14H long periodic stacking ordered structure precipitate, Scr. Mater. 53 (7) (2005) 799–803. [2] Y. Kawamura, T. Kasahara, S. Izumi, M. Yamasaki, Elevated temperature Mg97Y2Cu1 alloy with long period ordered structure, Scr. Mater. 55 (5) (2006) 453–456. [3] Y. Kawamura, M. Yamasaki, Formation and mechanical properties of Mg97Zn1RE2 alloys with long-period stacking ordered structure, Mater. Trans. 48 (11) (2007) 2986–2992. [4] T. Itoi, K. Takahashi, H. Moriyama, M. Hirohashi, A high-strength Mg–Ni–Y alloy sheet with a long-period ordered phase prepared by hot-rolling, Scr. Mater. 59 (10) (2008) 1155–1158. [5] S. Yoshimoto, M. Yamasaki, Y. Kawamura, Microstructure and mechanical properties of extruded Mg-Zn-Y alloys with 14H long period ordered structure, Mater. Trans. 47 (4) (2006) 959–965. [6] T. Itoi, T. Inazawa, M. Yamasaki, Y. Kawamura, M. Hirohashi, Microstructure and mechanical properties of Mg–Zn–Y alloy sheet prepared by hot-rolling, Mater. Sci. Eng. A 560 (2013) 216–223. [7] X.H. Shao, Z.Q. Yang, X.L. Ma, Strengthening and toughening mechanisms in Mg–Zn–Y alloy with a long period stacking ordered structure, Acta Mater. 58 (14) (2010) 4760–4771. [8] Z.Z. Peng, X.H. Shao, Q.Q. Jin, J.F. Liu, X.L. Ma, Dislocation configuration and solute redistribution of low angle kink boundaries in an extruded Mg–Zn–Y–Zr alloy, Mater. Sci. Eng. A 687 (2017) 211–220. [9] M. Matsuda, S. Ii, Y. Kawamura, Y. Ikuhara, M. Nishida, Variation of long-period
[12]
[13] [14] [15] [16]
[17] [18] [19]
[20] [21]
561
Precipitation No Mg24Y5 Mg24Y5
stacking order structures in rapidly solidified Mg97Zn1Y2 alloy, Mater. Sci. Eng. A 393 (1–2) (2005) 269–274. Y.M. Zhu, M. Weyland, A.J. Morton, K. Oh-ishi, K. Hono, J.F. Nie, The building block of long-period structures in Mg–RE–Zn alloys, Scr. Mater. 60 (11) (2009) 980–983. X.H. Shao, Z.Q. Yang, X.L. Ma, Interplay between deformation twins and basal stacking faults enriched with Zn/Y in Mg97Zn1Y2 alloy, Philos. Mag. Lett. 94 (3) (2014) 150–156. X.H. Shao, S.J. Zheng, D. Chen, Q.Q. Jin, Z.Z. Peng, X.L. Ma, Deformation twinning induced decomposition of lamellar LPSO structure and its re-precipitation in an MgZn-Y alloy, Sci. Rep. 6 (2016) 30096. W.W. Hu, Z.Q. Yang, H.Q. Ye, < c+a > dislocations and their interactions with other crystal defects in a Mg alloy, Acta Mater. 124 (2017) 372–382. S.-B. Mi, Q.-Q. Jin, New polytypes of long-period stacking ordered structures in Mg–Co–Y alloys, Scr. Mater. 68 (8) (2013) 635–638. Q.Q. Jin, X.H. Shao, X.B. Hu, Z.Z. Peng, X.L. Ma, New polytypes of LPSO structures in an Mg–Co–Y alloy, Philos. Mag. 97 (1) (2016) 1–16. Q.Q. Jin, X.H. Shao, Z.Z. Peng, J.H. You, K.Q. Qiu, X.L. Ma, Crystallographic account of an ultra-long period stacking ordered phase in an Mg88Co5Y7 alloy, J. Alloys Compd. 693 (2017) 1035–1038. Q.-Q. Jin, S.-B. Mi, Intermetallic phases in Mg–Co–Y alloys, J. Alloys Compd. 582 (2014) 130–134. M. Egami, E. Abe, Structure of a novel Mg-rich complex compound in Mg–Co–Y ternary alloys, Scr. Mater. 98 (2015) 64–67. Z. Yang, M.F. Chisholm, G. Duscher, X. Ma, S.J. Pennycook, Direct observation of dislocation dissociation and Suzuki segregation in a Mg–Zn–Y alloy by aberrationcorrected scanning transmission electron microscopy, Acta Mater. 61 (1) (2013) 350–359. X. Shao, Characterization of Long Period Stacking Ordered Phases and Their Effects on Mechanical Properties of Mg Alloys, Insititute of Metal Research Doctor, 2011. B.-J. Lv, J. Peng, Y. Peng, A.-T. Tang, F.-S. Pan, The effect of LPSO phase on hot deformation behavior and dynamic recrystallization evolution of Mg–2.0Zn–0.3Zr–5.8Y alloy, Mater. Sci. Eng. A 579 (2013) 209–216.
Materials Characterization 151 (2019) 553–562
Z.Z. Peng, et al.
[29] K. Kishida, Y. Takahama, H. Inui, c-axis compression twinning in an off-stoichiometric compound Ti3Al with the D019 structure, Mater. Sci. Eng. A 400-401 (2005) 339–344. [30] M.-X. Zhang, P.M. Kelly, Morphology and crystallography of Mg24Y5 precipitate in Mg–Y alloy, Scr. Mater. 48 (2003) 379–384. [31] D. Mizer, B.C. Perters, A study of precipitation at elevated temperatures in a Mg-8.7 pct Y alloy, Metall. Trans. 3 (1972) 3262–3264. [32] Z.Z. Peng, X.H. Shao, Q.Q. Jin, C.H. Li, X.L. Ma, Deformation induced precipitation of nano-particles in an Mg88Co5Y7 alloy, Scr. Mater. 116 (2016) 57–61. [33] N. Stanford, R. Cottam, B. Davis, J. Robson, Evaluating the effect of yttrium as a solute strengthener in magnesium using in situ neutron diffraction, Acta Mater. 78 (2014) 1–13. [34] X. Yang, Y. Jiang, Morphology and crystallographic characteristics of deformation bands in Mg alloy under hot deformation, Acta Metall. Sin. 46 (4) (2010) 451–457. [35] L.X. Wang, G. Fang, M.A. Leeflang, J. Duszczyk, J. Zhou, Constitutive behavior and microstructure evolution of the as-extruded AE21 magnesium alloy during hot compression testing, J. Alloys Compd. 622 (2015) 121–129. [36] B.-J. Lv, J. Peng, Y.-J. Wang, X.-Q. An, L.-P. Zhong, A.-T. Tang, F.-S. Pan, Dynamic recrystallization behavior and hot workability of Mg–2.0Zn–0.3Zr–0.9Y alloy by using hot compression test, Mater. Des. 53 (2014) 357–365.
[22] B.-J. Lv, J. Peng, L.-L. Zhu, Y.-J. Wang, A.-T. Tang, The effect of 14H LPSO phase on dynamic recrystallization behavior and hot workability of Mg–2.0Zn–0.3Zr–5.8Y alloy, Mater. Sci. Eng. A 599 (2014) 150–159. [23] X. Xia, Q. Chen, J. Li, D. Shu, C. Hu, S. Huang, Z. Zhao, Characterization of hot deformation behavior of as-extruded Mg–Gd–Y–Zn–Zr alloy, J. Alloys Compd. 610 (2014) 203–211. [24] N. Tahreen, D.F. Zhang, F.S. Pan, X.Q. Jiang, C. Li, D.Y. Li, D.L. Chen, Characterization of hot deformation behavior of an extruded Mg–Zn–Mn–Y alloy containing LPSO phase, J. Alloys Compd. 644 (2015) 814–823. [25] J.W. Christian, S. Mahajan, Deformation twinning, Prog. Mater. Sci. 39 (1–2) (1995) 1–157. [26] K. Máthis, K. Nyilas, A. Axt, I. Dragomir-Cernatescu, T. Ungár, P. Lukáč, The evolution of non-basal dislocations as a function of deformation temperature in pure magnesium determined by X-ray diffraction, Acta Mater. 52 (10) (2004) 2889–2894. [27] P. Perez, M. Eddahbi, S. Gonzalez, G. Garces, P. Adeva, Refinement of the microstructure during superplastic deformation of extruded Mg94Ni3Y1.5CeMM1.5 alloy, Scr. Mater. 64 (1) (2011) 33–36. [28] Z.J. Lin, M.J. Zhuo, Z.Q. Sun, P. Veyssiere, Y.C. Zhou, Amorphization by dislocation accumulation in shear bands, Acta Mater. 57 (9) (2009) 2851–2857.
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