- Email: [email protected]
Phase transformation of Mg–Li alloys induced by super-high pressure
Materials Letters 117 (2014) 45–48
Contents lists available at ScienceDirect
Materials Letters journal homepage: www.elsevier.com/locate/matlet
Phase transformation of Mg–Li alloys induced by super-high pressure Wenshi Wu, Qiuming Peng n, Jianxin Guo, Shuangshuang Zhao State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao 066004, PR China
art ic l e i nf o
a b s t r a c t
Article history: Received 19 October 2013 Accepted 23 November 2013 Available online 3 December 2013
The effect of superhigh pressure (4 GPa) on phase transformation behavior and mechanical properties of a duplex phase Mg–7 wt%Li alloy has been investigated in the temperature range from 450 to 1150 1C. Hardness and tensile strength of Mg–7 wt%Li alloy have been significantly improved by super-high pressure treatment at 700 1C for 120 min. The microstructural observation is confirmed that super-high pressure is of benefit for phase transformations from bcc-Li3Mg7 to hcp phases and hcp-Li3Mg17 to hcpLi0.92Mg4.08. The improved mechanical properties are mainly associated with the reduction of bcc-Li3Mg7 phase and the formation of {1 0 1 1} compression twins in hcp-Li0.92Mg4.08 phase. & 2013 Elsevier B.V. All rights reserved.
Keywords: Phase transformation Mechanical properties Mg–Li alloys
1. Introduction Mg–Li based alloys exhibit attractive applications in aerospace and aircraft because of their super stiffness to weight ratio and excellent cold formability. Depending on Li concentration, Mg–Li system can be basically divided into three structures, i.e. hcp-α, hcpαþbcc-β and bcc-β. Wherein, the duplex phase (hcp-αþ bcc-β) alloy containing 5.7–11 wt% of Li is highly desirable for structural materials, which can provide an optimum combination of mechanical properties. However, its absolute strength is low, which hinders its wide applications [1]. Therefore, to date alloying with a third element such as Al, Zn, Si, Ag, Cd or rare earth elements, and grain refinement via thermo-mechanical processing are commonly used to improve strength of duplex Mg–Li alloy [2,3]. Unfortunately, both of them are very expensive or complex to obtain homogeneous and high strength Mg–Li alloys. It is a well-known fact that exterior pressure plays an important role in phase formation [4]. More recently, it is reported that super-high pressure synthesis is one of the most effective techniques to prepare new compounds or to obtain unique properties [5]. Unlike precipitation strengthening and grain refinement reported previously, it is expected that SHP treatment would provide a simple method to improve mechanical properties of Mg–Li alloys by modifying phase composition and its morphology. Herein, a super-high pressure (SHP) approach is successfully introduced to improve the hardness and strength of Mg–7 wt%Li alloy for the first time. 2. Experimental process The ingot with a nominal composition of Mg–7Li (wt%, all compositions given thereafter in wt%) was prepared by vacuum n
Corresponding author. Tel.: þ 86 335 80570 47; fax: þ86 335 8074 545. E-mail address: [email protected] (Q. Peng).
0167-577X/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.matlet.2013.11.094
induction melting method. The ingot was then machined into samples with 10 mm in diameter and 8 mm in length for superhigh pressure (SHP) treatment. The samples were wrapped with Ta foil and then inserted into a BN crucible. The pressure (4 GPa) was loaded before heating the samples. The samples were heated to different temperatures from 450 to 1150 and held at that temperature for 120 min. Hereafter the samples were quenched to room temperature before unloading the pressure. All the SHPed samples were heated to 200 1C for 5 min to release residual stress. Microhardness (HV) was tested with a Vickers hardness tester. The load and the dwelling time were 100 g and 15 s, respectively. Comparatively, Rockwell hardness (RH) was also tested on the bulk SHPed samples with 1 mm ball indenter and a load of 10 kgf. Compression tests were performed using a Gleeble-3500 thermomechanical simulator with a strain rate of 1.7 10 3 s 1 at room temperature. X-ray diffraction (XRD) was carried out on Xpert-Pro diffractometer with Cu Kα radiation in the range from 201 to 801 with 0.21 min 1. The results were analyzed by Rietveld refinement. Microstructural observations were carried out on a JEM2010 transmission electron microscope (TEM) operated at 200 kV. TEM thin foils were prepared by ion milling in a Fischione system operating at 5 keV at an incidence angle ranging between 7 and 151.
3. Results and discussion Fig. 1a shows hardness variation of SHPed samples under different temperatures. Both HV and RH rapidly increase and then decrease with increasing processing temperature. The sample SHPed at 700 1C (hereafter denoted as 4 GPa-700 sample) shows the highest hardness, which is nearly two times higher than that of pristine sample. The representative true stress–strain compression curves are shown in Fig. 1b. The pristine Mg–7Li alloy shows very low yield strength while high uniform deformation.
46
W. Wu et al. / Materials Letters 117 (2014) 45–48
The sample was not ruptured until reaching to 100% strain. After SHPed at 700 1C, the yield strength and ultimate compressive strength were remarkably improved, and the values are 286 MPa and 311 MPa, respectively. Nevertheless, the rupture strain was decreased to 50%. When further increased the SHP temperature to 1050 1C, the rupture strain was recovered to 70%, and the alloy reveals the similar strengths as compared with the as-cast alloy. Fig. 2a shows XRD patterns of different state samples. The pristine Mg–7Li alloy is mostly composed of α and β phases. The β phase corresponds to bcc-Li3Mg7 (PDF#65-6742) intermetallic compound. However, the α phase contains two stoichiometric compounds, which are hcp-Li0.92Mg4.08 (PDF#65-4080) and hcpLi3Mg17 (PDF#65-5512). Fig. 2b shows a localized high magnified image of XRD pattern, it can be identified that the (1 0 2) peak of both hcp phases is shifted, and intensity changes greatly under
different SHP temperatures. Notably, the (1 0 2) peak intensity of hcp-Li3Mg17 is higher than that of hcp-Li0.92Mg4.08 in addition to 4 GPa-700 sample. Depending on XRD results, the phase relative abundance (PRA) of alloys is analysed by Rietveld software (Fig. 2c). It can be seen from Fig. 2d that the PRA value of bcc-Li3Mg7 phase in pristine sample is 0.27, and it reduces to 0.06 in 4 GPa-700 sample. With further increasing the temperature, the PRA will gradually increase, which is ascribed to a re-solidification process. Briefly, when the temperature is higher than the melting point, the SHP treatment becomes a single solidification under 4 GPa. Therefore, when the temperature is higher than 700 1C, the alloys show the similar morphology and phase composition as pristine one. Regarding to hcp phases, the PRA of hcp-Li0.92Mg4.08 and hcpLi3Mg17 compounds was changed significantly when the SHP temperature is increased. For comparison, the volume ratio (λ)
Fig. 1. (a) Effect of superhigh pressure on microhardness and Rockwell hardness; (b) representative compressive curves of pristine sample, 4GPa-700 and 4GPa-1050 samples at room temperature.
Fig. 2. (a) X-ray patterns of different state Mg–7Li samples; (b) local high magnification of X-ray curves in (a); (c) Rietveld analysis pattern for the pristine sample; (d) effect of SHP on the phase relative abundance of bcc-Li3Mg7 and the ratio of volume fraction (Li0.92Mg4.08/Li3Mg17).
W. Wu et al. / Materials Letters 117 (2014) 45–48
47
Fig. 3. (a) A typical TEM micrograph of pristine sample; (b) SAED of A1 phase along [ 1 1 0 1] in (a); (c) SAED of B1 phase along [0 0 0 1] in (a); (d) a typical TEM micrograph of the 4 GPa-700 sample; (e) SAED of A2 phase along [ 1 1 0 1] in (d); (f) SAED of B2 phase along [ 1 1 0 1] in (d); (g) schematic representation of orientation relationship of compression twin in (f).
for the hcp phases is introduced [6]: λ ¼ δLi0:92 Mg4:08 =δLi3 Mg17 , where δ is a volume fraction. The λ also exhibits the convex trend (Fig. 2d), which is consistent with the HV and RH results. The value of pristine sample is 0.21, whilst the maximum value of 1.85 is observed in 4 GPa-700 sample. Fig. 3a shows TEM bright-field micrograph of as-cast sample. The selected area electron diffraction (SAED) patterns taken from A1 and B1 areas are shown in Fig. 3b and Fig. 3c, respectively. Both A1 and B1 reveal typical hcp structure, which correspond to [ 1 1 0 1] and [0 0 0 1] zone axis, respectively. The lattice parameters of A1 phase are a ¼3.182 Å and c ¼5.090 Å, whilst those of B1 phase are a ¼3.132 Å and c ¼5.071 Å. The C phase is confirmed as bcc-Li3Mg7 compound (not shown). In the case of 4 GPa-700 sample, the α phase mostly consists of two lamellar phases (A2 and B2). The width of B2 phase is measured to be 7207 20 nm, which is 3 times thicker than that of A2 phase. The SAED patterns of the lamellar phases are shown in Fig. 3e and f, respectively. Both of them are taken from the [ 1 1 0 1] direction. The a and c lattices of A2 phase is 3.074 Å and 4.932 Å, whilst the values are 3.041 Å and 4.902 Å for B2 phase. According to the XRD results and previous literature [7,8], it can be confirmed that the A1 and A2 phases correspond to hcp-Li3Mg17 phase. The B1 and B2 phases are hcp-Li0.92Mg4.08 phase. In addition, as shown in Fig. 3d and f, some twins are also observed in the B2 phase. The schematic diagram of twin orientation relationship is shown in Fig. 3g. A {1 0 1 1} compression twin along 561 〈 1 0 1 2〉 was identified in the B2 phase. Combining with microstructural analysis, it is considered that the improvement of mechanical properties by SHP treatment is possibly associated with the decrease of bcc-Li3Mg7 phase, the phase transformation from hcp-Li3Mg17 to hcp-Li0.92Mg4.08 and the formation of {1 0 1 1} compression twin in hcp-Li0.92Mg4.08 phase.
Phase transformation induced the improvement of mechanical properties has been commonly observed in age hardenable Mg alloys. However, unlike conventional aging treatment, the SHP treatment studied in this work is performed under super-high pressure and relatively high temperature. Levitas [9] pointed out that under pressure, the dislocations tend to pile up at strong barriers owing to the difference in yield stress between two phases, giving rise to very high local stresses. In turn, it can enhance the phase transformation. The promotion of phase transformation under SHP has been confirmed in carbon [10], CaB4 [11] etc. Herein, under the high pressure (4 GPa), the dislocations pile up at the interface of soft/hard phase such as bcc-Li3Mg7/hcp-Li3Mg17 or bccLi3Mg7/hcp-Li0.92Mg4.08. With increasing the temperature, the dislocations will easily cut through the softer bcc phases, leading to its partial dissolution. Therefore, the volume fraction of β phase is reduced and the volume fraction of α phase is increased. Similar phase transformation is confirmed in a single phase bcc-type Mg–Li alloy under high pressure torsion [2]. In addition, SHP treatment will lead the phase transformation from hcp-Li3Mg17 to hcp-Li0.92Mg4.08. However, taking into account the compressibilityðκÞ, the hardness is proportional to the bulk modulus (B) [12,13]: HV p mB ¼ a1κ , where m and a are constants. According to the first principle calculations [14], the bulk modulus of hcp-Li0.92Mg4.08 ( 30 GPa) is lower than that of hcp-Li3Mg17 compound ( 35 GPa), indicating that the increase of hcp-Li0.92Mg4.08 phase will lead to the decrease of strength. However, it is believed that the formation of superfine twin will improve the mechanical properties of Mg–Li alloys. According to Estrin's model [15], the spacing between impenetrable obstacles or grain boundaries decides the mean free path of dislocations, in which the mean free path is determined both by dislocation– dislocation interaction and by such barriers as grain and twin
48
W. Wu et al. / Materials Letters 117 (2014) 45–48
Fig. 4. TEM micrographs of 4 GPa–700 Mg–7Li alloy with a strain of 1%. (a) some strain areas are observed in grain boundaries; (b) local high magnification of (a), where dislocation aggregation is detected in both twin boundaries (TB) and interior of twin (IT).
boundaries. Since twin formation subdivides grain as twin boundaries, the formation of twin offers additional barriers to dislocation movement. As shown in Fig. 4a, a large number of deformed areas are observed in grain boundaries. Under higher magnification (Fig. 4b), the tangled dislocations are observed in both twin boundaries (TB) and interior of twin (IT), which demonstrates that the presence of compression twin provides an effective role in prohibiting dislocation movement during the deformation. 4. Conclusions The super high pressure treatment is an effective approach to improve the mechanical properties of duplex phase Mg–Li alloys. The high pressure can significantly induce phase transformation. The improved hardness and strength of Mg–7 wt%Li alloy after high pressure is mostly attributed to the decrease of bcc-Li3Mg7 phase and the presence of compression twin in hcp-Li0.92Mg4.08 phase. Acknowledgment The contributions of Dr Wenlong Xiao in postprocessing are acknowledged. This research is supported by NSFC (51101142,
50821001 and 51102206), New Century Excellent Talents in University of Ministry of Education of China (NCET-12-0690) and Heibei province scientific program (13961002D and Y2012019).
References [1] Gasior W, Moser Z, Zakulski W, Schwitzgebel G. Metall Mater Trans A 1996;27:2419. [2] Srinivasarao B, Zhilyaev AP, Gutiérrez-Urrutia I, Pérez-Pradoa MT. Scr Mater 2013;68:583. [3] Liu T, Wang YD, Wu SD, Lin Peng R, Huang CX, Jiang CB, et al. Scr Mater 2004;51:1057. [4] Tian Y, Xu B, Yu D, Ma Y, Wang Y, Jiang Y, et al. Nature 2013;493:385. [5] Jie JC, Zou CM, Wang HW, Li B, Wei ZJ. Scr Mater 2011;64:588. [6] Young RA, Wiles DB. J Appl Crystallogr 1982;15:430. [7] Hardie D, Parkins RN. Philos Mag 1959;4:815. [8] Herbstein FH, Averbach BL. Acta Metall 1956;4:407. [9] Levitas VI. J Mech Phys Solids 1997;45:923. [10] Li Q, Ma Y, Oganov AR, Wang H, Wang H, Xu Y, et al. Phys Rev Lett 2009;102:175506. [11] Liu Z, Han X, Yu D, Sun Y, Xu B, Zhou X-F, et al. Appl Phys Lett 2010;96:031903. [12] Léger JM, Haines J. Endeavour 1997;21:121. [13] Léger JM, Haines J, Blanzat B. J Mater Sci Lett 1994;13:1688. [14] Peng Q, Wu W, Guo J, Xiang J, Zhao S., J Appl Phys Accepted. [15] Estrin Y, Mecking H. Acta Metall 1984;32:57.
Contents lists available at ScienceDirect
Materials Letters journal homepage: www.elsevier.com/locate/matlet
Phase transformation of Mg–Li alloys induced by super-high pressure Wenshi Wu, Qiuming Peng n, Jianxin Guo, Shuangshuang Zhao State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao 066004, PR China
art ic l e i nf o
a b s t r a c t
Article history: Received 19 October 2013 Accepted 23 November 2013 Available online 3 December 2013
The effect of superhigh pressure (4 GPa) on phase transformation behavior and mechanical properties of a duplex phase Mg–7 wt%Li alloy has been investigated in the temperature range from 450 to 1150 1C. Hardness and tensile strength of Mg–7 wt%Li alloy have been significantly improved by super-high pressure treatment at 700 1C for 120 min. The microstructural observation is confirmed that super-high pressure is of benefit for phase transformations from bcc-Li3Mg7 to hcp phases and hcp-Li3Mg17 to hcpLi0.92Mg4.08. The improved mechanical properties are mainly associated with the reduction of bcc-Li3Mg7 phase and the formation of {1 0 1 1} compression twins in hcp-Li0.92Mg4.08 phase. & 2013 Elsevier B.V. All rights reserved.
Keywords: Phase transformation Mechanical properties Mg–Li alloys
1. Introduction Mg–Li based alloys exhibit attractive applications in aerospace and aircraft because of their super stiffness to weight ratio and excellent cold formability. Depending on Li concentration, Mg–Li system can be basically divided into three structures, i.e. hcp-α, hcpαþbcc-β and bcc-β. Wherein, the duplex phase (hcp-αþ bcc-β) alloy containing 5.7–11 wt% of Li is highly desirable for structural materials, which can provide an optimum combination of mechanical properties. However, its absolute strength is low, which hinders its wide applications [1]. Therefore, to date alloying with a third element such as Al, Zn, Si, Ag, Cd or rare earth elements, and grain refinement via thermo-mechanical processing are commonly used to improve strength of duplex Mg–Li alloy [2,3]. Unfortunately, both of them are very expensive or complex to obtain homogeneous and high strength Mg–Li alloys. It is a well-known fact that exterior pressure plays an important role in phase formation [4]. More recently, it is reported that super-high pressure synthesis is one of the most effective techniques to prepare new compounds or to obtain unique properties [5]. Unlike precipitation strengthening and grain refinement reported previously, it is expected that SHP treatment would provide a simple method to improve mechanical properties of Mg–Li alloys by modifying phase composition and its morphology. Herein, a super-high pressure (SHP) approach is successfully introduced to improve the hardness and strength of Mg–7 wt%Li alloy for the first time. 2. Experimental process The ingot with a nominal composition of Mg–7Li (wt%, all compositions given thereafter in wt%) was prepared by vacuum n
Corresponding author. Tel.: þ 86 335 80570 47; fax: þ86 335 8074 545. E-mail address: [email protected] (Q. Peng).
0167-577X/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.matlet.2013.11.094
induction melting method. The ingot was then machined into samples with 10 mm in diameter and 8 mm in length for superhigh pressure (SHP) treatment. The samples were wrapped with Ta foil and then inserted into a BN crucible. The pressure (4 GPa) was loaded before heating the samples. The samples were heated to different temperatures from 450 to 1150 and held at that temperature for 120 min. Hereafter the samples were quenched to room temperature before unloading the pressure. All the SHPed samples were heated to 200 1C for 5 min to release residual stress. Microhardness (HV) was tested with a Vickers hardness tester. The load and the dwelling time were 100 g and 15 s, respectively. Comparatively, Rockwell hardness (RH) was also tested on the bulk SHPed samples with 1 mm ball indenter and a load of 10 kgf. Compression tests were performed using a Gleeble-3500 thermomechanical simulator with a strain rate of 1.7 10 3 s 1 at room temperature. X-ray diffraction (XRD) was carried out on Xpert-Pro diffractometer with Cu Kα radiation in the range from 201 to 801 with 0.21 min 1. The results were analyzed by Rietveld refinement. Microstructural observations were carried out on a JEM2010 transmission electron microscope (TEM) operated at 200 kV. TEM thin foils were prepared by ion milling in a Fischione system operating at 5 keV at an incidence angle ranging between 7 and 151.
3. Results and discussion Fig. 1a shows hardness variation of SHPed samples under different temperatures. Both HV and RH rapidly increase and then decrease with increasing processing temperature. The sample SHPed at 700 1C (hereafter denoted as 4 GPa-700 sample) shows the highest hardness, which is nearly two times higher than that of pristine sample. The representative true stress–strain compression curves are shown in Fig. 1b. The pristine Mg–7Li alloy shows very low yield strength while high uniform deformation.
46
W. Wu et al. / Materials Letters 117 (2014) 45–48
The sample was not ruptured until reaching to 100% strain. After SHPed at 700 1C, the yield strength and ultimate compressive strength were remarkably improved, and the values are 286 MPa and 311 MPa, respectively. Nevertheless, the rupture strain was decreased to 50%. When further increased the SHP temperature to 1050 1C, the rupture strain was recovered to 70%, and the alloy reveals the similar strengths as compared with the as-cast alloy. Fig. 2a shows XRD patterns of different state samples. The pristine Mg–7Li alloy is mostly composed of α and β phases. The β phase corresponds to bcc-Li3Mg7 (PDF#65-6742) intermetallic compound. However, the α phase contains two stoichiometric compounds, which are hcp-Li0.92Mg4.08 (PDF#65-4080) and hcpLi3Mg17 (PDF#65-5512). Fig. 2b shows a localized high magnified image of XRD pattern, it can be identified that the (1 0 2) peak of both hcp phases is shifted, and intensity changes greatly under
different SHP temperatures. Notably, the (1 0 2) peak intensity of hcp-Li3Mg17 is higher than that of hcp-Li0.92Mg4.08 in addition to 4 GPa-700 sample. Depending on XRD results, the phase relative abundance (PRA) of alloys is analysed by Rietveld software (Fig. 2c). It can be seen from Fig. 2d that the PRA value of bcc-Li3Mg7 phase in pristine sample is 0.27, and it reduces to 0.06 in 4 GPa-700 sample. With further increasing the temperature, the PRA will gradually increase, which is ascribed to a re-solidification process. Briefly, when the temperature is higher than the melting point, the SHP treatment becomes a single solidification under 4 GPa. Therefore, when the temperature is higher than 700 1C, the alloys show the similar morphology and phase composition as pristine one. Regarding to hcp phases, the PRA of hcp-Li0.92Mg4.08 and hcpLi3Mg17 compounds was changed significantly when the SHP temperature is increased. For comparison, the volume ratio (λ)
Fig. 1. (a) Effect of superhigh pressure on microhardness and Rockwell hardness; (b) representative compressive curves of pristine sample, 4GPa-700 and 4GPa-1050 samples at room temperature.
Fig. 2. (a) X-ray patterns of different state Mg–7Li samples; (b) local high magnification of X-ray curves in (a); (c) Rietveld analysis pattern for the pristine sample; (d) effect of SHP on the phase relative abundance of bcc-Li3Mg7 and the ratio of volume fraction (Li0.92Mg4.08/Li3Mg17).
W. Wu et al. / Materials Letters 117 (2014) 45–48
47
Fig. 3. (a) A typical TEM micrograph of pristine sample; (b) SAED of A1 phase along [ 1 1 0 1] in (a); (c) SAED of B1 phase along [0 0 0 1] in (a); (d) a typical TEM micrograph of the 4 GPa-700 sample; (e) SAED of A2 phase along [ 1 1 0 1] in (d); (f) SAED of B2 phase along [ 1 1 0 1] in (d); (g) schematic representation of orientation relationship of compression twin in (f).
for the hcp phases is introduced [6]: λ ¼ δLi0:92 Mg4:08 =δLi3 Mg17 , where δ is a volume fraction. The λ also exhibits the convex trend (Fig. 2d), which is consistent with the HV and RH results. The value of pristine sample is 0.21, whilst the maximum value of 1.85 is observed in 4 GPa-700 sample. Fig. 3a shows TEM bright-field micrograph of as-cast sample. The selected area electron diffraction (SAED) patterns taken from A1 and B1 areas are shown in Fig. 3b and Fig. 3c, respectively. Both A1 and B1 reveal typical hcp structure, which correspond to [ 1 1 0 1] and [0 0 0 1] zone axis, respectively. The lattice parameters of A1 phase are a ¼3.182 Å and c ¼5.090 Å, whilst those of B1 phase are a ¼3.132 Å and c ¼5.071 Å. The C phase is confirmed as bcc-Li3Mg7 compound (not shown). In the case of 4 GPa-700 sample, the α phase mostly consists of two lamellar phases (A2 and B2). The width of B2 phase is measured to be 7207 20 nm, which is 3 times thicker than that of A2 phase. The SAED patterns of the lamellar phases are shown in Fig. 3e and f, respectively. Both of them are taken from the [ 1 1 0 1] direction. The a and c lattices of A2 phase is 3.074 Å and 4.932 Å, whilst the values are 3.041 Å and 4.902 Å for B2 phase. According to the XRD results and previous literature [7,8], it can be confirmed that the A1 and A2 phases correspond to hcp-Li3Mg17 phase. The B1 and B2 phases are hcp-Li0.92Mg4.08 phase. In addition, as shown in Fig. 3d and f, some twins are also observed in the B2 phase. The schematic diagram of twin orientation relationship is shown in Fig. 3g. A {1 0 1 1} compression twin along 561 〈 1 0 1 2〉 was identified in the B2 phase. Combining with microstructural analysis, it is considered that the improvement of mechanical properties by SHP treatment is possibly associated with the decrease of bcc-Li3Mg7 phase, the phase transformation from hcp-Li3Mg17 to hcp-Li0.92Mg4.08 and the formation of {1 0 1 1} compression twin in hcp-Li0.92Mg4.08 phase.
Phase transformation induced the improvement of mechanical properties has been commonly observed in age hardenable Mg alloys. However, unlike conventional aging treatment, the SHP treatment studied in this work is performed under super-high pressure and relatively high temperature. Levitas [9] pointed out that under pressure, the dislocations tend to pile up at strong barriers owing to the difference in yield stress between two phases, giving rise to very high local stresses. In turn, it can enhance the phase transformation. The promotion of phase transformation under SHP has been confirmed in carbon [10], CaB4 [11] etc. Herein, under the high pressure (4 GPa), the dislocations pile up at the interface of soft/hard phase such as bcc-Li3Mg7/hcp-Li3Mg17 or bccLi3Mg7/hcp-Li0.92Mg4.08. With increasing the temperature, the dislocations will easily cut through the softer bcc phases, leading to its partial dissolution. Therefore, the volume fraction of β phase is reduced and the volume fraction of α phase is increased. Similar phase transformation is confirmed in a single phase bcc-type Mg–Li alloy under high pressure torsion [2]. In addition, SHP treatment will lead the phase transformation from hcp-Li3Mg17 to hcp-Li0.92Mg4.08. However, taking into account the compressibilityðκÞ, the hardness is proportional to the bulk modulus (B) [12,13]: HV p mB ¼ a1κ , where m and a are constants. According to the first principle calculations [14], the bulk modulus of hcp-Li0.92Mg4.08 ( 30 GPa) is lower than that of hcp-Li3Mg17 compound ( 35 GPa), indicating that the increase of hcp-Li0.92Mg4.08 phase will lead to the decrease of strength. However, it is believed that the formation of superfine twin will improve the mechanical properties of Mg–Li alloys. According to Estrin's model [15], the spacing between impenetrable obstacles or grain boundaries decides the mean free path of dislocations, in which the mean free path is determined both by dislocation– dislocation interaction and by such barriers as grain and twin
48
W. Wu et al. / Materials Letters 117 (2014) 45–48
Fig. 4. TEM micrographs of 4 GPa–700 Mg–7Li alloy with a strain of 1%. (a) some strain areas are observed in grain boundaries; (b) local high magnification of (a), where dislocation aggregation is detected in both twin boundaries (TB) and interior of twin (IT).
boundaries. Since twin formation subdivides grain as twin boundaries, the formation of twin offers additional barriers to dislocation movement. As shown in Fig. 4a, a large number of deformed areas are observed in grain boundaries. Under higher magnification (Fig. 4b), the tangled dislocations are observed in both twin boundaries (TB) and interior of twin (IT), which demonstrates that the presence of compression twin provides an effective role in prohibiting dislocation movement during the deformation. 4. Conclusions The super high pressure treatment is an effective approach to improve the mechanical properties of duplex phase Mg–Li alloys. The high pressure can significantly induce phase transformation. The improved hardness and strength of Mg–7 wt%Li alloy after high pressure is mostly attributed to the decrease of bcc-Li3Mg7 phase and the presence of compression twin in hcp-Li0.92Mg4.08 phase. Acknowledgment The contributions of Dr Wenlong Xiao in postprocessing are acknowledged. This research is supported by NSFC (51101142,
50821001 and 51102206), New Century Excellent Talents in University of Ministry of Education of China (NCET-12-0690) and Heibei province scientific program (13961002D and Y2012019).
References [1] Gasior W, Moser Z, Zakulski W, Schwitzgebel G. Metall Mater Trans A 1996;27:2419. [2] Srinivasarao B, Zhilyaev AP, Gutiérrez-Urrutia I, Pérez-Pradoa MT. Scr Mater 2013;68:583. [3] Liu T, Wang YD, Wu SD, Lin Peng R, Huang CX, Jiang CB, et al. Scr Mater 2004;51:1057. [4] Tian Y, Xu B, Yu D, Ma Y, Wang Y, Jiang Y, et al. Nature 2013;493:385. [5] Jie JC, Zou CM, Wang HW, Li B, Wei ZJ. Scr Mater 2011;64:588. [6] Young RA, Wiles DB. J Appl Crystallogr 1982;15:430. [7] Hardie D, Parkins RN. Philos Mag 1959;4:815. [8] Herbstein FH, Averbach BL. Acta Metall 1956;4:407. [9] Levitas VI. J Mech Phys Solids 1997;45:923. [10] Li Q, Ma Y, Oganov AR, Wang H, Wang H, Xu Y, et al. Phys Rev Lett 2009;102:175506. [11] Liu Z, Han X, Yu D, Sun Y, Xu B, Zhou X-F, et al. Appl Phys Lett 2010;96:031903. [12] Léger JM, Haines J. Endeavour 1997;21:121. [13] Léger JM, Haines J, Blanzat B. J Mater Sci Lett 1994;13:1688. [14] Peng Q, Wu W, Guo J, Xiang J, Zhao S., J Appl Phys Accepted. [15] Estrin Y, Mecking H. Acta Metall 1984;32:57.