ZHANG, S., SUN, Y., WU, R., WANG, X., CHEN, X.-B., FERNANDEZ, C., and PENG, Q. 2020. Coherent interface strengthening of ultrahigh pressure heat-treated Mg-Li-Y alloys. Journal of materials science and technology [online], 51, pages 79-83. Available from: https://doi.org/10.1016/j.jmst.2020.02.039

Coherent interface strengthening of ultrahigh pressure heat-treated Mg-Li-Y alloys.

ZHANG, S., SUN, Y., WU, R., WANG, X., CHEN, X.-B., FERNANDEZ, C., and PENG, Q.

2020

This document was downloaded from https://openair.rgu.ac.uk Journal of Materials Science & Technology

CoherentFor Interface Strengthening of Ultrahigh Pressure Heat-Treated Mg-Li-Y Alloys Journal: JournalRevie of Materials Science & Technology Manuscript ID JMST-2019-3402.R1

Manuscript Type: Letter

Date Submitted by the 27-Dec-2019 Author: w Complete List of Authors: Zhang, Shun; Yanshan University, State Key Laboratory of Metastable Materials Science and Technology; Engineering University, Key Laboratory of Superlight Materials & Surface Technology, Ministry of Education Sun, Yong; Yanshan University,Only State Key Laboratory of Metastable Materials Science and Technology Wu, Ruizhi; Harbin Engineering University, Key Laboratory of Superlight Materials & Surface Technology, Ministry of Education Wang, Xiang; Harbin Engineering University, Key Laboratory of Superlight Materials & Surface Technology, Ministry of Education Chen, Xiao-Bo; RMIT University, School of Engineering Fernandez, Carlos; Robert Gordon University, School of Pharmacy and life sciences Peng, Qiuming; Yanshan University, State Key Laboratory of Metastable Materials Science and Technology

Keywords: Mg alloys, Nano-twins, Stacking faults, Strengthening

Speciality: Magnesium Alloys

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1 2 3 4 Letter 5 6 7 Coherent Interface Strengthening of Ultrahigh Pressure Heat-Treated 8 9 Mg-Li-Y Alloys 10 11 Shun Zhang 1,2, Yong Sun 1, Ruizhi Wu 2, Xiang Wang *,2, Xiao-Bo Chen 3, Carlos Fernandez 4, 12 13 *,1 14 Qiuming Peng 15 1State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, 16 17 , 066004, 18 19 2Key Laboratory of Superlight Materials & Surface Technology, Ministry of Education, Harbin 20 21 Engineering University, Harbin, 150001, China 22 23 3School of Engineering, RMIT University, Carlton 3053, VIC, Australia 24 25 4School of Pharmacy and life sciences, Robert Gordon University, Aberdeen, AB107GJ, UK 26 27 28 * Corresponding author. Prof. Qiuming Peng; Prof. Xiang Wang 29 30 31 E-mail address: [email protected] (Qiuming Peng); [email protected] (Xiang 32 33 Wang) 34 35 36 37 38 Achieving good strength-ductility of Mg alloys has always been a crucial issue for the 39 40 widespread applications of Mg-based structural materials. Herein, an unexpected 41 42 double-stage strengthening phenomenon was discovered in Mg-8Li-1Y (wt.%) alloys 43 44 through high-pressure (6 GPa) heat treatments over a range of 700-1300 °C. Attractively, 45 46 the yield strength values are improved remarkably without losing their ductility. The 47 low-temperature strengthening mechanism is mainly driven by the formation of 48 49 50 large-volume nanoscale contraction twins. In contrast, the high-temperature 51 52 strengthening reason is ascribed to the presence of densely nano-sized stacking faults. 53 54 Both coherent interfaces contribute effectively to high mechanical strength without any 55 56 tradeoff in ductility. 57 58 59 60 Key words: Mg alloys; Nano-twins; Stacking faults; Strengthening 1

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1 2 3 4 1. Introduction 5 6 7 8 Mg alloys, the lightest structural metals, play an important role in energy saving and emission 9 10 reduction [1]. However, due to the low strength and poor formability at room temperature, their 11 12 engineering applications are still unsatisfied. A large number of approaches have been attempted to 13 14 improve the strength of Mg alloys in the past decades [2]. Traditional strengthening strategies, such 15 16 as grain refinement, precipitation and sub-grain boundaries through dislocation plugging, could 17 18 introduce a large number of boundaries into Mg matrix, which can effectively prevent dislocation 19 20 motion [2]. However, it is common to see the increment in mechanical strength at the expense of 21 22 ductility through those incoherent interfaces, which is ascribed to the limited capacity to 23 24 accommodate dislocations [3]. Hence, the feasible approach to attain high strength without 25 26 scarification in plasticity remains a challenge for the wide applications of Mg alloys. 27 28 Differing from incoherent boundaries, coherent twin boundaries (TBs), can not only act as 29 30 effective barriers to dislocation movement, but also provide huge space for dislocation storage. This 31 32 strategy is applicable for FCC metals [4] whilst difficult for HCP Mg, owing to the intrinsic 33 characteristics of twins in Mg alloys. In particular, the interfaces of [1012] tension twins (TTWs) 34 _

35 are not stable. In addition, though the interfaces of [10 1] contraction twins (CTWs) are more stable 36 1 37 than those of TTWs, secondary TTWs are readily initiated in and occupy the entire internal domains 38 39 of CTWs [5]. Therefore, it is difficult to strengthen Mg alloys by increasing _ volume fractions of 40 _

41 TBs. More recently, our group introduced distinctive hierarchical {1011}−{1011} double 42 43 contraction nanotwins in Mg-8Li (all compositions given thereafter in wt.%) binary alloy by means 44 of an ultra-high pressure treatment (UHP) technique [6]. Such a unique structure elevates the yield 45 46 47 strength of Mg-8Li alloys up from 50 to 249 MPa while maintains its high ductility (23.6% 48 49 elongation rate). Subsequently, an improved contraction nanotwins-stacking faults structure was 50 51 achieved by increasing Li concentration to 13 wt.%, which further increases the specific strength of 52 53 Mg-13Li alloy [7]. It is anticipated that coherent structure strengthening could be a feasible strategy 54 to design high strength-ductile Mg alloys. 55 56 In the present work, we attempt to explore the effects of Y alloying element on interface 57 58 strengthening of Mg-8Li alloys. It is well-known that large-sized rare earth elements (REs) are 59 60 prone to segregate in TBs, which remarkably affects their structures and properties [8]. 2

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1 2 3 4 Experimental work was designed to provide evidence to elucidate two critical issues that dominate 5 coherent interface strengthening in Mg alloys: i) is it technically feasible to yield coherent structure 6 7 8 strengthening through REs alloying under milder conditions? and ii) which structure of contraction 9 10 twins or stacking faults is more effective to improve mechanical properties? 11 12 13 2. Experimental processes 14 15 16 17 A Mg-8Li-0.5Zn-1Y (LZY801) ingot was prepared by vacuum induction melting and machined 18 into samples with dimensions of 10 mm in diameter and 14 mm in length for following 19 20 21 UHPed treatments. Samples were wrapped up with Ta foil and then inserted into a boron nitride 22 23 crucible in a cubic-anvil large-volume press with six rams. Loading pressure of 6 GPa was operated 24 before heating up the system. After UHPed treatment holding at a defined temperature (over a range of 25 26 700 -1300 °C) for 30 min, LZY801 specimens were quenched down to room temperature. The 27 sample was then heated from room temperature to 200 ° C at a heating rate of 5 ° C / min to release 28 29 30 residual stress. LZY801 samples which were subject to UHPed treatment under 6 GPa at 900 °C for 30 31 32 min was denoted as LZY801-6GPa-900. The as-cast counterpart was introduced as a reference. 33 34 Microhardness was measured through a Vickers hardness tester (HV) with an applied load of 100 35 36 g and dwelling time of 15 s. Mechanical tests were performed on Gleeble 3500 with a strain rate for 37 38 tension and compression tests of 0.4 mm/min and 1×10-3 s-1, respectively. Length and width of the 39 40 dog-bone shape tensile samples were 6 and 1.5 mm, respectively. In comparison, compressive samples 41 42 were accurately machined into dimensions of 5 mm in diameter and 10 mm in length. Microstructural 43 44 investigations were performed using optical and transmission electron microscopy (TEM). TEM thin 45 46 foils were prepared by Ar ion milling in a Fischione system operating at 5 keV. X-ray diffraction 47 48 (XRD) was carried out on a Rigatu 2500C with Cu-Kα radiation at a scanning rate of 4°/min over a 49 50 range of 20 - 80°. 51 52 53 3. Results and discussion 54 55 56 57 Figure 1a illustrates microhardness evolution of the UHPed samples as a function of processing 58 59 temperature. The as-cast LZY801sample presents the lowest microhardness value of 48 60 3

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1 2 3 4 5 6 7 8 HV, similar to that of previous as-cast Mg-8Li alloy [6]. It is noted that the samples are 9 dramatically hardened (75 HV) from the original as-cast state as a result of heat treatment 10 11 (700-1300 °C) at a constant high pressure (6 GPa). Two distinct strengthening stages are evident 12 13 between 700-900 °C and 1000-1200 °C, respectively, followed by a sharp drop in microhardness 14 15 from its peak point of 96 HV (1200 °C) to 54 HV (1300 °C). 16 In general, the melting point of a metal increases with externally applied pressure, such as the 17 18 representative cases of Al-Zn alloys [9]. As such, two processes, i.e. ultra-high solid solution and 19 ultra-high re-solidification, would occur during the UHPed heat treatments. Namely, the UHPed 20 21 solid solution process occurs in a range of 700-1200 °C, whilst the UHPed re-solidification is 22 triggered above 1200 °C, which is confirmed by both optical microstructure and XRD patterns 23 24 (Figure 1b). Differing from α phase (Mg-rich) and β phase (Li-rich), the eutectic YZn5 (hexagonal 25 with space group P6/mmm) phase is reduced with increasing temperature, whilst the equal-axis grain 26 27 structure well remains below 1200 °C. Analogous to the atmospheric heat treatment, the UPed 28 treatment mainly involves two typical processes. If the temperature is below melting point, the heat 29 30 treatment is related to the UPed solid solution. Comparatively, when the temperature is over melting 31 point, the heat treatment is ascribed to the UPed re-solidification. According to our previous results 32 33 [6,7], the melting point of Mg alloys will be increased to over 1200 ˚C under 6GPa. As soon as the 34 temperature is over melting point, the strengthening structure is eliminated, and the main 35 36 strenthening role is merely related to solid-solution strengthening. Therefore, is demonstrated that 37 the increment in microhardness of UHPed LZY801 is related to the UHPed solid solution treatment, 38 39 while the decrease of microhardness is related to the ultra-high re-solidification. 40 41

42 43 Figure 1c displays the representative true stress-strain curves of the as-cast, the 44 45 LZY801-6GPa-900 and the LZY801-6GPa-1200 specimens, which agrees with the results of 46 47 microhardness measurements (Figure 1a). A low yield strength of 89 MPa in the as-cast LZY801 48 49 alloy is largely increased up to 248 MPa through heat treatment at 900 °C (LZY801-6GPa-900), 50 51 and approaches the maximum value of 270 MPa for the LZY801-6GPa-1200, three times of that of 52 53 the as-cast one. It is noteworthy that such improvement in strength does not sacrifice elongation in 54 55 a significant manner (~19 %, Figure 1c). Same results are confirmed by compressive tests, 56 57 where samples are not fractured until it reaches the load limitation of compression tester (Figure 58 59 1d). In addition, the tensile-compression asymmetry (TCA)-the ratio of tensile yield 60 strength 4

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1 2 3 4 (σtys)/compressive yield strength (σcys)-is significantly minimized after UHPed treatment, and the TCA value is reduced from 1.25 (as-cast) to 1.02 (LZY801-6GPa-1200). In particular, tensile yield 65 7 8 strength of all LZY801-6GPa alloys exceeds largely those of both Mg-Li binary alloys (100 ± 50 9 MPa) and majority of multiple Mg-Li-X alloy systems (X: other alloying elements), even after 10 11 severe plastic deformation treatments, such as rolling [10-14], extrusion [14-17], equal- 12 13 channel angular pressing [18, 19] and aging [20]. 14 15 TEM characterization was employed to elucidate the mechanisms of the two-stage 16 17 strengthening. Regarding the as-cast case (Figure 3a-c), both α and β phases are evidently 18 19 distributed in the matrix. Moreover, a number of particles with a size of 0.1 μm in diameter are 20 21 accumulated in grain boundaries. Corresponding energy dispersive spectroscopy (EDS) mapping 22 23 (Figure 3b) confirms that the precipitates in the as-cast LZY801 sample are enriched in Y. 24 25 Therefore, those particles are assigned to eutectic YZn5 phase, as indicated in the XRD pattern 26 27 (Figure 1b). In the case of the LZY801-6GPa-900 (Figure 3d), a few nanoscale twins are 28 29 embedded within the matrix, and some nanoscale particles are found. HRTEM micrograph (Figure 30 31 3e) exhibits that nano-twin boundaries (NTBs) are mainly composed of coherent structures, and the 32 33 angle between (0001) twin and (0001) matrix is ~123.4 °, indicating the formation of CTWs [5]. The 34 35 average twin thickness is about 80 nm. 36 37 Figure 3f-g reveals that the dimension of the particles is 10-20 nm. The spacing of (100) is 38 39 about 0.453 nm, which aligns will with the theoretical value (0.432 nm) of YZn5 [21]. Consequently, 40 41 it is evident that the strengthening stage I is attributed to the presence of such a unique 42 microstructure, i.e. co-presence of a large number of nano-coherent twins in matrix and a 43 few residual nanoscale YZn 44 5 phases. According to the empirical twin-strengthening mode [6], 45 46 47 theoretical yield strength with NTWs can be calculated by the following equation: 48 -1/2 49 ∆σy = nkTBλ (1) 50 1/2 51 Where λ is 80 nm; n is the volume fraction; and kTB is the Hall-Petch constant (0.35 MPa m ). 52 53 Hence, the theoretical increment in yield strength of CTWs is 1.23 GPa. 54 In contrast, a large number of SFs with an average spacing of 18 nm exist in the matrix of 55 56 LZY801-6GPa-1200 alloy without any twins or YZn particles (Figure 4a). The formation of SFs is 57 5 58 also confirmed by the streaks in SAED (Figure 4b). The atomic-scale image of SFs (Figure 4c) 59 60 indicates that the SFs are assigned to growth fault (I1-type), differing from the stacking sequence 5

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1 2 3 4 (...ABABABAB...) of an ideal HCP Mg alloy. It is normally built by removing one A plane above 5 6 one B plane, and then shearing the remaining planes above B plane by 1/3[1100], resulting in 7 a …ABACBCBCB… structure. According to previous SFs strengthening mechanisms, the 8 9 10 theoretical increased yield strength can be calculated by the following equation [7]: 11 -1 12 ∆σy = mkSFd (2) 13 14 Where d is the average width of SFs (18 nm); m is the volume fraction, and kSF is experimental 15 constant (3780 MPa nm). Therefore, the theoretical increase in yield strength of SFs is 2.6 GPa. 16 17 18 4. Conclusions 19 20 21 The effect of Y alloying element on mechanical properties of the UHPed Mg-8Li alloys was 22 23 explored. The results reveal that the addition alloying element can obviously optimize the UHPed 24 25 conditions by reducing the operating temperature. Moreover, it demonstrates that both CTWs and 26 27 SFs are achieved within one simple system by varying the UHPed conditions, and SFs are more 28 29 effective to the enhancement in mechanical properties. The tradeoff between strength and ductility 30 of Mg alloys can be mitigated through modifying nanoscale coherent interfaces. 31 32 33 34 35 Acknowledgments 36 37 38 We greatly acknowledge the financial support from NSFC (51771162 and 51971194) and 39 40 province talent foundation (A201910002). 41 42 43 References 44 45 [1] X.J. Wang, D.K. Xu, R.Z. Wu, X.B. Chen, Q.M. Peng, L. Jin, Y.C. Xin, Z.Q. Zhang, Y. Liu, 46 47 X.H. Chen, G. Chen, K.K. Deng, H.Y. Wang, J. Mater. Sci. Technol. 34 (2018) 245-247. 48 49 [2] T.M. Pollock, Science. 328 (2010) 986-987. 50 51 [3] W. Xu, N. Birbilis, G. Sha, Y. Wang, J.E. Daniels, Y. Xiao, M. Ferry, Nat. Mater. 14 (2015) 52 53 1229. 54 55 [4] X.Y. Li, Y.J. Wei, L. Lu, K. Lu, H.J. Gao, Nature. 464 (2010) 877-880. 56 57 [5] P. Cizek, M.R. Barnett, Scr. Mater. 59 (2008) 959-962. 58 59 [6] H. Fu, B.C. Ge, Y.C. Xin, R.Z. Wu, C. Fernandez, J.Y. Huang, Q.M. Peng, Nano. Lett. 17 60 (2017) 6117-6124. 6

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1 2 3 4 5 6 [7] Q.M. Peng, Y. Sun, B.C. Ge, H. Fu, Q. Zu, X.Z. Tang, J.Y. Huang, Acta. Mater. 169 (2019) 7 8 36-44. 9 [8] J.F. Nie, Y.M. Zhu, J.Z. Liu, X.Y. Fang, Science. 340 (2013) 957-960. 10 11 [9] B. Ahn, A.P. Zhilyaev, H.-J. Lee, M. Kawasaki, T.G. Langdon, Mat. Sci. Eng. A 635 (2015) 12 13 109-117. 14 15 [10] C.J. Shuai, W.J. Yang, Y.W. Yang, C.D. Gao, C.X. He, H. Pan, Int. J. Bioprint. 5 (2019) 49. 16 17 [11] G.Y. Sha, X.G. Sun, T. Liu, Y.H. Zhu, T. Yu, J. Mater. Sci. Technol. 27 (2011) 753-758. 18 19 [12] D.K. Xu, T.T. Zu, M. Yin, Y.B. Xu, E.H. Han, J. Alloys. Compd. 582 (2014) 161-166. 20 21 [13] T.L. Zhu, J.F. Sun, C.L. Cui, R.Z. Wu, S. Betsofen, Z. Leng, J.H. Zhang, M.L. Zhang, Mat. Sci. 22 23 Eng. A 600 (2014) 1-7. 24 25 [14] M.L. Zhang, R.Z. Wu, T. Wang, J. Mater. Sci. 44 (2009) 1237-1240. 26 27 [15] C.Q. Li, D.K. Xu, B.J. Wang, L.Y. Sheng, R.Z. Wu, E.H. Han, J. Mater. Sci. Technol. 35 28 29 (2019) 2477-2484. 30 31 [16] Y. Tang, W.T. Jia, X. Liu, Q.C. Le, Y.L. Zhang, Mat. Sci. Eng. A 675 (2016) 55-64. 32 33 [17] Y. Yang, X.M. Xiong, J.F. Su, X.D. Peng, H.M. Wen, G.B. Wei, F.S. Pan, E.J. Lavernia, J. 34 35 Alloys. Compd. 750 (2018) 696-705. 36 37 [18] M. Karami, R. Mahmudi, Mat. Sci. Eng. A 607 (2014) 512-520. 38 39 [19] R.Z. Wu, Z.K. Qu, M.L. Zhang, Rev. Adv. Mater. Sci. 24 (2010) 35-43. 40 41 [20] C.O. Muga, H. Guo, S.S. Xu, Z.W. Zhang, Mat. Sci. Eng. A 689 (2017) 195-202. 42 43 [21] M. Fornasini, J. Less-Common. Met. 25 (1971) 329-332. 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 7

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1 2 3 4 Figure captions 5 Figure 1. Mechanical properties of UHPed LZY801 alloys as a function of processing temperature. 6 7 8 (a) Effects of UHP conditions on microhardness. (b) Optical microstructure and XRD patterns. (c) 9 10 Tensile and (d) compression curves. The insets correspond to macro-size optical photographs. 11 12 13 14 Figure 2. Comparison in terms of mechanical properties between UHPed LZY801 alloys and a 15 16 range of representative Mg-Li-X (alloying elements) alloys prepared by a variety of 17 themal-mechanical treatments. 18 19 20 21 Figure 3. TEM characterization of the LYZ801-6GPa-900 alloy. (a) STEM micrographs of the 22 23 as-cast LZY801 sample. (b) EDS mapping of Y in (a). (c) SAED pattern of α-Mg. (d) TEM image 24 25 of the LZY801-6GPa-900 observed along [1120] axis. A randomly particle in the matrix is marked 26 27 by a red arrow. (e) Atomic-level image of coherent contraction twin boundary. (f) High 28 29 magnification image of the particle in (d). (g) Local atomic-level magnified region as marked in (f). 30 31 32 33 Figure 4. Strengthening mechanisms in the LZY801-6GPa-1200 alloy. (a) TEM image with an 34 35 inset corresponding to the spacing distribution SFs derived from statistical calculations of 50 36 37 random results. (b) Corresponding SAED patterns, where crystal planes and SFs (streaking pattern) 38 39 are marked. (c) Atomic level image of SFs in (a). 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 8

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1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 Figure 1. Mechanical properties of UHPed LZY801 alloys as a function of processing 34 35 36 temperature. (a) Effects of UHP conditions on microhardness. (b) Optical microstructure and 37 38 39 XRD patterns. (c) Tensile and (d) compression curves. The insets correspond to macro-size 40 41 optical photographs. 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 9

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1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 Figure 2. Comparison in terms of mechanical properties between UHPed LZY801 alloys and a 25 26 range of representative Mg-Li-X (alloying elements) alloys prepared by a variety of 27 28 29 themal-mechanical treatments. 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 10

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1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 Figure 3. TEM characterization of the LYZ801-6GPa-900 alloy. (a) STEM micrographs of 36 37 38 the as-cast LZY801 sample. (b) EDS mapping of Y in (a). (c) SAED pattern of α-Mg. (d) 39 40 41 TEM image of the LZY801-6GPa-900 observed along [1120] axis. A randomly particle in the 42 43 matrix is marked by a red arrow. (e) Atomic-level image of coherent contraction twin 44 45 46 boundary. (f) High magnification image of the particle in (d). (g) Local atomic-level 47 48 magnified region as marked in (f). 49 50 51 52 53 54 55 56 57 58 59 60 11

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1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 Figure 4. Strengthening mechanisms in the LZY801-6GPa-1200 alloy. (a) TEM image with an 36 37 38 inset corresponding to the spacing distribution SFs derived from statistical calculations of 50 39 40 random results. (b) Corresponding SAED patterns, where crystal planes and SFs (streaking 41 42 43 pattern) are marked. (c) Atomic level image of SFs in (a). 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 12

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1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 For 21 22 23 24 Revie 25 26 27 28 29 w 30 31 Figure 1. Mechanical properties of UHPed LZY801 alloys as a function of processing temperature. (a) Effects 32 of UHP conditions on microhardness. (b) Optical microstructure and XRD patterns. (c) Tensile and (d) 33 compression curves. The insets correspond to macro-size optical photographs. 34 35 Only 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 https://mc03.manuscriptcentral.com/jmst Journal of Materials Science & Technology Page 14 of 16

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 For 21 22 23 24 Revie 25 26 27 28 29 Figure 2. Comparison in terms of wmechanical properties between UHPed LZY801 alloys and a range of 30 representative Mg-Li-X (alloying elements) alloys prepared by a variety of themal-mechanical treatments. 31 32 33 34 Only 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 https://mc03.manuscriptcentral.com/jmst Page 15 of 16 Journal of Materials Science & Technology

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 For 21 22 23 24 Revie 25 26 27 28 29 w 30 31 32 33 34 Only 35 36 37 Figure 3. TEM characterization of the LYZ801-6GPa-900 alloy. (a) STEM micrographs of the as-cast LZY801 38 sample. (b) EDS mapping of Y in (a). (c) SAED pattern of α-Mg. (d) TEM image of the LZY801-6GPa-900 39 observed along [11"2" 0] axis. A randomly particle in the matrix is marked by a red arrow. (e) Atomic-level 40 image of coherent contraction twin boundary. (f) High magnification image of the particle in (d). (g) Local 41 atomic-level magnified region as marked in (f). 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 https://mc03.manuscriptcentral.com/jmst Journal of Materials Science & Technology Page 16 of 16

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 For 21 22 23 24 Revie 25 26 27 28 29 w 30 31 32 33 34 Only 35 36 37 38 39 40 Figure 4. Strengthening mechanisms in the LZY801-6GPa-1200 alloy. (a) TEM image with an inset 41 corresponding to the spacing distribution SFs derived from statistical calculations of 50 random results. (b) 42 Corresponding SAED patterns, where crystal planes and SFs (streaking pattern) are marked. (c) Atomic 43 level image of SFs in (a). 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60 https://mc03.manuscriptcentral.com/jmst