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Re-recognition of the aging precipitation behavior in the Mg–Sm binary alloy
Journal of Alloys and Compounds 814 (2020) 152320
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Re-recognition of the aging precipitation behavior in the MgeSm binary alloy Hongbo Xie a, Boshu Liu a, Junyuan Bai a, Changli Guan a, Dongfang Lou a, Xueyong Pang a, Hong Zhao a, Shanshan Li a, Yuping Ren a, b, *, Hucheng Pan a, **, Changlin Yang c, Gaowu Qin a, b a
Key Laboratory for Anisotropy and Texture of Materials (Ministry of Education), School of Materials Science and Engineering, Northeastern University, Shenyang, 110819, China Research Center for Metallic Wires, Northeastern University, Shenyang, 110819, China c State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi'an, 710072, China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 5 June 2019 Received in revised form 10 September 2019 Accepted 16 September 2019 Available online 17 September 2019
Generally, the MgeSm system is one of the magnesium light rare-earth alloys (Mg-LRE), and its aging precipitation behavior is similar to that of the MgeNd or MgeCe alloys. For a long time, the b phase, precipitated after the b1 intermediate phase, has been considered to be Mg12Sm, although this has never been experimentally verified. In this study, however, we clearly confirm that the b phase is Mg5Sm with a face-centered cubic structure (space group: F43m; a ¼ 22.34 Å) by aberration-corrected scanning transmission electron microscopy (STEM) combined with quantitative STEM simulation. This is similar to the case of magnesium heavy rare-earth alloys (Mg-HRE). The lenticular-shaped b plates have a {1010}a habit plane and is fully coherent with the a-Mg matrix, with the orientation relationships among the b phase, b1 phase, and a-Mg matrix being (111)b//(111)b1//(1120)a and [110]b//[110]b1//[0001]a. This finding improves our understanding of the precipitation/strengthening mechanism in the MgeSm based alloys and is expected to guide the future design of novel high-strength heat-resistant Mg-RE based alloys. © 2019 Elsevier B.V. All rights reserved.
Keywords: Magnesium alloys Aging Precipitation Crystal structure HAADF-STEM Simulation
1. Introduction Magnesium alloys have long been considered as one of the potential weight-saving light structural materials [1e12]. However, their relatively poor mechanical properties, low ductility, and poor heat-resistant behaviors restrict their widespread applications [1e5]. Recently, Mg rare-earth (RE) alloys have received considerable attention due to their high strength at either room or elevated temperatures, excellent creep properties, and good formability [9,10]. The improvement of their mechanical properties is mainly ascribed to their high-density precipitates on the non-basal planes. However, the precipitation behaviors of these alloys are complex and closely related to those of the RE types [9,13]; the structures of the intermediate precipitates change with time at high
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (Y. Ren), [email protected] (H. Pan). https://doi.org/10.1016/j.jallcom.2019.152320 0925-8388/© 2019 Elsevier B.V. All rights reserved.
temperatures, resulting in the deterioration of their mechanical properties during service. In general, Mg-RE alloys can be classified simply into two groups based on the difference in the atomic numbers and chemical properties of the RE elements [9,13]: (i) Mg-LRE alloys, including Mg-Ce [14,15], Mg-Nd [16e19], and MgeSm [20e23]. The precipitation sequence of the Mg-LRE alloys is commonly agreed to be in the order of supersaturated solid solution (S$S$S$S) / b’’ (D019, Mg3RE) / bS’ (bco, Mg7RE) / b1 (fcc, Mg3RE) / b (bct, Mg12RE) [9,13e21]. (ii) Mg-HRE alloys, including Mg-Gd [24e30], Mg-Dy [31], and MgeY [25,29,32]. The precipitation sequence of the MgHRE alloys is generally accepted to be in the order of S$S$S$S / b’’ (D019, Mg3RE) / bL’ (bco, Mg7RE) / b1 (fcc, Mg3RE) / b (fcc, Mg5RE/cubic, Mg41RE5) [9,24e32]. From the precipitation sequences, it can be seen that one of the noticeable differences between the Mg-LRE and Mg-HRE alloys is the structure of the b0 intermediate phase, which is the main strengthening precipitate phase on the {1010}a plane. For the MgLRE alloys, the b0 precipitate is determined to have the short-
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structure b’ (denoted as bS’), having a composition of Mg7RE and an ordered orthorhombic structure (a ¼ 6.42 Å, b ¼ 11.12 Å, c ¼ 5.21 Å) [9,16e23]. For the Mg-HRE alloys, in contrast, the b0 phase has the long-structure (denoted as bL’), also with a composition of Mg7RE but an ordered orthorhombic structure (a ¼ 6.42 Å, b ¼ 22.23 Å, c ¼ 5.21 Å) [9,25e27,29e32], and the difference is the lattice parameter of b. However, notably, recent studies have indicated that both the bS’ and bL’ precipitates could co-exist in either MgeSm or MgeGd binary alloys [23,26,27]. Another difference in the aging precipitate sequence is the structure of the b precipitate after the b1 phase. For the Mg-HRE alloys, such as the MgeGd system, the b phase has been identified as the Mg5Gd phase with a face-centered cubic structure [9,24,28]. For the Mg-LRE alloys, however, all the phases precipitated after the b1 are thought to be the rod-shaped Mg12RE phase, having a body-centered tetragonal structure, such as Mg12Ce, Mg12Nd, and Mg12Sm [9,23]. Unfortunately, the b precipitate of Mg12Sm has not been confirmed by any transmission electron microscopy (TEM) or high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) studies to date. Therefore, this study aims to use aberration-corrected HAADFSTEM, selected area electron diffraction (SAED), and quantitative STEM simulation to characterize the b phase precipitated in a binary Mg-3wt.%Sm alloy during isothermal aging. Additionally, we unambiguously confirm that the b phase is a new Mg5Sm intermetallic compound (fcc, space group: F43m; a ¼ 22.34 Å), which is an unusual phenomenon in Mg-RE alloys that until now, had been thought unlikely. This finding improved our understanding of the precipitation/strengthening mechanism in MgeSm based alloys, and it could facilitate relevant modeling and simulation studies for this alloy system.
2. Experimental procedures The alloy with nominal composition of Mg-3wt.%Sm (Mg0.5 at.%Sm) was prepared by melting the pure Mg (99.9 wt%) and Mg-25Sm (wt.%) master alloy in the induction furnace with protection of the argon atmosphere. The molten alloy was stirred and kept at 760 C for 5 min and poured into a steel mold preheated to 300 C. The chemical composition of the obtained as-cast ingot was measured by OPTIMA 4300 DV composition analyzer, and the actual compositions was determined to be Mg-2.84Sm (wt.%). The
as-cast samples were solution treated at 520 C for 12 h, followed by water quenching and ageing in oil bath at 200 C for 300 h, and finally continued isothermally aged at 300 C for 48 h. The TEM specimens with a diameter of 3 mm were prepared by twin jet electro-polishing at - 40 C in mixture solution of 5.3 g lithium chloride, 11.2 g magnesium perchlorate, 500 ml methanol and 100 ml 2-butoxy ethanol, and subsequently ion milling with low energy electron beam. Finally, Gatan SOLARUS (950) Plasma Cleaning System was used to clean up the sample surfaces. TEM and HAADF-STEM observations were carried out using the JEMARM200F at an accelerating voltage of 200 kV, equipped with probe Cs corrector and cold field emission gun. The probe convergence is 25 mrad which yields a probe size of less than 0.1 nm, and the camera length was set to 8 cm which yields a collection semiangle of 48e327 mrad. The present study employed the quantitative TEM/STEM simulations software to simulate HAADF-STEM images, and the corresponding parameters used for simulations are identical with the experimental settings: a voltage of 200 kV, the third order spherical aberration C3 ¼ 0.05 mm, defocus values df ¼ 13.7 nm, beam convergence angle a ¼ 21.4 mrad.
3. Results and discussions Fig. 1a shows a low-magnification HAADF-STEM image recorded from the Mg-3Sm alloy isothermally aged at 200 C for 300 h, viewed along the [0001]a direction. It can be seen that there are high-density precipitates with high contrast in the a-Mg matrix, and these lenticular-shaped precipitates with a length more than ~200 nm distribute uniformly in three directions: ½1120a, ½1210a, and ½2110a. Furthermore, the majority of the precipitates in this aging-stage were identified as the b1 phase (Mg3Sm, fcc, space group: Fm3m; a ¼ 0.74 nm), which was reported in many Mg-RE based alloys [9,33e36]. Noteworthily, the orientation relationships between the b1 phase and the a-Mg matrix are [110]b1// [0001]a; (111)b1//(1120)a [9,33,35]. The atomic-scale HAADF-STEM image of one precipitate in Fig. 1a is enlarged and shown in Fig. 1b, and the b-Mg5Sm phase clearly forms in the b1 intermediate precipitate (as shown by the cyan circle). This observation result is consistent with the results of a recent study in the binary MgeGd and the WE54 alloys [28,34], a similar orientation relationships exists between the b-Mg5RE phase and the a-Mg matrix, i.e., [110]b//[0001]a; (111)b//(1120)a [9,24,28,34]. Additionally, the
Fig. 1. HAADF-STEM images of the Mg-3Sm alloy isothermally aged at 200 C for 300 h. (a) Low-magnification HAADF-STEM image. (b) A bright-contrast precipitate in Fig. 1a is enlarged and shown in Fig. 1b; the atomic-scale HAADF-STEM image indicated that there are some b-Mg5Sm phases formed in the over-aged state. The electron beam is parallel to the [0001]a direction.
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lattice image of the b phase is marked with a yellow rhombus in Fig. 1b, and the inset is its corresponding modeled atomic arrangement. At present, it is generally accepted that the b1 phase forms prior to the b phase, as evidenced by literature [9,24,28,34]. Our experimental observation indicates that the b phase nucleates in the existing b1 precipitate and grows gradually by the inside-out mode. It is generally believed that the Mg5RE phase can only precipitate in Mg-HRE alloys (MgeGd) during isothermal aging, and the terminal phase of the aging precipitation in binary MgeSm alloys is considered to be the Mg12Sm phase, instead of Mg5Sm. In order to characterize the detailed structure of the Mg5Sm phase observed above, and to confirm its thermal stability and evolution, the agedsample above was further isothermally aged at 300 C for 48 h. The TEM bright-field images and SAED patterns of the Mg-3Sm alloy further isothermally aged at 300 C for 48 h, viewed along the [0001]a, and [1120]a zone axes, respectively, are shown in Fig. 2. The lenticular-shaped precipitates with low contrast are observed to distribute uniformly in the ½1120a, ½1210a, and ½2110a directions, and coarsen up to a length of more than 500 nm and a thickness of ~100 nm. The corresponding SAED patterns, as shown in Fig. 2c and d, exhibit extra diffraction spots (labeled by yellow numbers) between the a-Mg matrix diffraction spots (labeled by red numbers). This indicates that the precipitates have an fcc structure, and it can be identified as the Mg5Sm phase (fcc, space group: F43m; a ¼ 22.34 Å), which is consistent with the Mg5Gd
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equilibrium phase [28,34]. Besides, from the superimposed precipitate and matrix SAED patterns it can be determined definitely that the orientation relationships between the b phase and the aMg matrix are (111)b//(1120)a and [110]b//[0001]a. Furthermore, this implies that the orientation relationships among the b phase, b1 phase, and a-Mg matrix are (111)b//(111)b1//(1120)a and [110]b// [110]b1//[0001]a. Additional details concerning the b phase were obtained by conducting the HAADF-STEM analysis. Fig. 3a provides a lowmagnification HAADF-STEM image of the Mg-3Sm alloy further isothermally aged at 300 C for 48 h, viewed along the [0001]a direction, showing the b phase with a {1010}a habit plane (high contrast) precipitated in the a-Mg matrix, which is consistent with the bright-field TEM image shown in Fig. 2a. Fig. 3b shows two b precipitates connected at an angle of 120 , and the connection area, marked by a cyan circle in Fig. 3b, is further enlarged and presented in Fig. 3c. It can be seen that the b/b interface is incoherent. Fig. 3d presents a continuous b precipitate containing two kinds of contrasts, and the connection area marked by the cyan circle in Fig. 3d is further enlarged and shown in Fig. 3e. The local part of a b precipitate is further enlarged, as shown in Fig. 3f, and the atomic-scale HAADF-STEM image shows that there are many bright dots inside the b precipitate. Each bright dot represents a column rich in Sm atoms because the brightness of individual atomic columns in the HAADF-STEM images approximates to the square of the average atomic numbers (the atomic numbers for Mg and Sm are 12 and 62,
Fig. 2. Bright-field TEM images and the corresponding SAED patterns of the Mg-3Sm alloy with continued isothermal aging at 300 C for 48 h. (a and b) Bright-field TEM images. (c and d) Corresponding SAED patterns, and the cyan circles in Fig. 2a and b correspond to the SAED areas. The electron beam is parallel to the [0001]a (a and c) and [1120]a (b and d) directions.
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Fig. 3. HAADF-STEM images of the Mg-3Sm alloy subjected to continuous isothermal aging at 300 C for 48 h. (a) Low-magnification HAADF-STEM image. (b and c) HAADF-STEM image in Fig. 3b show two b precipitates connected an angle of 120 , and the connection area marked by a cyan circle is further enlarged and presented in Fig. 3c. (d and e) HAADFSTEM image in Fig. 3d shows a b precipitate containing two kinds of contrasts, and the connection area marked by a cyan circle is further enlarged and shown in Fig. 3e. (f) Atomicscale HAADF-STEM image. The electron beam is parallel to the [0001]a direction.
respectively) [37e39]. It can be seen that a periodic unit-cell with a 70.5 rhombus structure is marked by yellow dotted lines in this image, and the careful measurement of the HAADF-STEM image indicates that the value of the lattice parameter, a, is ~22.34 Å. The inset in Fig. 3f is the modeled atomic arrangement for the b precipitate viewed along the [110]b direction. In this model, Mg atoms are shown with red spheres and Sm atoms displayed with blue spheres, and it can be seen that the marked Sm atom arrangement shows perfect conformity with the HAADF-STEM results, both in symmetry and in interatomic distance. This verifies that the b phase precipitated in the binary MgeSm alloys is the Mg5Sm intermetallic compound. The Mg-3Sm aged alloy was further analyzed by HAADF-STEM viewed along the ½1120a direction, as shown in Fig. 4. Fig. 4a provides a low-magnification HAADF-STEM image, and the local part of this image, marked by a cyan square frame, is further enlarged and shown in the inset; the perfect match of the (110)b/ (0001)a planes indicates that the relationship of the b/a-Mg interface is fully coherent. The atomic-scale HAADF-STEM image is shown in Fig. 4b, in which a periodic lattice with an equilateral hexagon structure and a side length of ~9.12 Å is marked by yellow dotted lines; the inset in Fig. 4b shows that the ½111b modeled atomic arrangement of the b-Mg5Sm phase is in excellent agreement with the experimental observations above. Fig. 5 presents the HAADF-STEM simulation results for the Mg5Sm phase. The three-dimensional view of the unit cell is shown in Fig. 5a. It can be seen that the Mg5Sm phase has an fcc structure with a side length of 22.34 Å, and the corresponding simulated HAADF-STEM results viewed along the [110]b and ½111b directions are shown in Fig. 5b and c, respectively. The unique characteristics of the b phase, i.e., the rhombus lattice images (in the [110]b direction) and equilateral hexagon lattice images (in the ½111b direction) that consist of the bright dots, are clearly visible in the
simulated HAADF-STEM results, which is well consistent with the experimental HAADF-STEM observations shown in Figs. 3 and 4. This indicates that the b-Mg5Sm phase can be generated and stably exist. The Mg12RE phases, including the b-Mg12Nd, was originally regarded as the equilibrium phase but is, in fact, a metastable phase, and the equilibrium precipitate phase is be-Mg41Nd5 [9,40e42]. This is consistent with the MgeNd binary phase diagram. According to the MgeSm binary phase diagram, the equilibrium phase should be Mg41Sm5 (space group: I4/m; a ¼ 14.74 Å, c ¼ 10.40 Å); however, similar to the Mg12Sm phase, the Mg41Sm5 phase has never been confirmed by any electron diffraction or TEM techniques in the aging precipitation sequence. The age-hardenable MgeSm based alloy has relatively high mechanical properties attributed to its high-density b0 and b1 intermediate precipitates on the non-basal planes of the a-Mg matrix; however, its precipitation sequence is different from the previous one [20e23]. Since the precipitation behavior in the earlyor peak-aged stages is similar to those of the MgeCe and MgeNd alloys, its over-aged precipitation behavior was considered to be consistent with the Mg-LRE alloys, i.e., the Mg12Sm precipitate formed after the b1 phase. However, we unambiguously confirmed that the b phase precipitated in the over-aged MgeSm binary alloys is Mg5Sm, a new phase with a structure similar to that of Mg5Gd, not Mg12Sm. In other words, the over-aged precipitation behavior of the MgeSm alloys is consistent with that of the Mg-HRE alloys [9,24e30]. Additionally, the bS’ and bL’ intermediate precipitates have been confirmed to co-exist in the MgeSm alloys [23]. Therefore, from the viewpoint of the precipitation behavior, the MgeSm system should be classified as a Mg middle rare-earth alloy between the Mg-LRE and Mg-HRE alloys. Conversely, although the b’ and b1 intermediate phases have significant strengthening effects in the MgeSm alloys, their easy coarsening and change into the b
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Fig. 4. HAADF-STEM images of the Mg-3Sm alloy subjected to continuous isothermal aging at 300 C for 48 h. (a) Low-magnification HAADF-STEM image. (b) Atomic-scale HAADFSTEM image. The electron beam is parallel to the [1120]a direction.
precipitation in the MgeSm binary alloy, was investigated in this work. The b phase is the Mg5Sm intermetallic compound (fcc, space group: F43m; a ¼ 22.34 Å), which has a structure that is consistent with that of Mg5Gd, not that of the presently accepted Mg12RE. The lenticular-shaped b plates with a {1010}a habit plane distribute uniformly in three directions: ½1120a, ½1210a, and ½2110a, which is fully coherent with the a-Mg matrix. The orientation relationships among the b phase, b1 phase, and a-Mg matrix are (111)b//(111)b1// (1120)a and [110]b//[110]b1//[0001]a. In addition, the precipitation sequence of the MgeSm binary alloys should be modified as follows: S$S$S$S / intermediate phases / b1 (fcc, Mg3Sm) / b (fcc, Mg5Sm). The b1 phase has poor thermal stability at 300 C. This finding is important for understanding the precipitation mechanism of the MgeSm binary alloy, and is expected to guide the future design of novel high-strength MgeSm based alloys with high thermal stability. Acknowledgements The authors are grateful to the financial support from the National Key Research and Development Program of China (Grant No. 2016YFB0701202), the National Natural Science Foundation of China (Grant No. 51371046, No. 51525101, No. 51501032, No. U1610253), the Fundamental Research Funds of the Central Universities (Grant No. N141008001, No. N170204011), the Project of Promoting Talents in Liaoning Province (No. XLYC1808038), and the fund of the state Key Laboratory of Solidification Processing in NPU (Grant No. SKLSP201920). The experiments were conducted at the Institute of Advanced Material Technology of Northeastern University and we acknowledge the use of the aberration-corrected transmission electron microscopy. Fig. 5. HAADF-STEM simulation results of the Mg5Sm phase. (a) Three-dimensional view of the unit cell of the Mg5Sm phase. (b) Simulated [110]b HAADF-STEM image. (c) Simulated ½111b HAADF-STEM images.
phase would deteriorate the mechanical properties of the allow. Improving the thermal stability of these intermediate phases would be a challenging topic in the development of Mg-Sm-based alloys.
4. Conclusions The structure of the b phase, which is the terminal of the aging
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[24] X. Gao, S.M. He, X.Q. Zeng, L.M. Peng, W.J. Ding, J.F. Nie, Microstructure evolution in a Mg-15Gd-0.5Zr (wt.%) alloy during isothermal aging at 250 C, Mater. Sci. Eng. A 431 (2006) 322e327. [25] H. Liu, Y. Gao, J.Z. Liu, Y.M. Zhu, Y. Wang, J.F. Nie, A simulation study of the shape of b0 precipitates in Mg-Y and Mg-Gd alloys, Acta Mater. 61 (2013) 453e466. [26] H. Liu, W.F. Xu, N.C. Wilson, L.M. Peng, J.F. Nie, Formation of and interaction between bF0 and b0 phases in a Mg-Gd alloy, J. Alloy. Comp. 712 (2017) 334e344. [27] H.B. Xie, H.C. Pan, Y.P. Ren, S.N. Sun, L.Q. Wang, Y.F. He, G.W. Qin, Co-existences of the two types of b’ precipitations in peak-aged Mg-Gd binary alloy, J. Alloy. Comp. 738 (2018) 32e36. [28] H.B. Xie, H.C. Pan, Y.P. Ren, H.X. Li, M. Jiang, G.W. Qin, Atomic-scale characterization of the equilibrium b-Mg5Gd phase by means of HAADF-STEM, Chin. J. Stereol. Image Anal 24 (2019) 91e97. [29] H. Liu, Y. Gao, J.Z. Liu, Y.M. Zhu, Y. Wang, J.F. Nie, A simulation study of the shape of b0 precipitates in Mg-Y and Mg-Gd alloys, Acta Mater. 61 (2013) 453e466. [30] J.X. Zheng, X.S. Xu, K.Y. Zhang, B. Chen, Novel structures observed in Mg-GdY-Zr during isothermal ageing by atomic-scale HAADF-STEM, Mater. Lett. 152 (2015) 287e289. [31] K. Saito, A. Yasuhara, M. Nishijima, K. Hiraga, Structural changes of precipitates by aging of an Mg-4 at% Dy solid solution studied by atomic-scaled transmission electron microscopy, Mater. Trans. 52 (2011) 1009e1015. [32] M. Nishijima, K. Yubuta, K. Hiraga, Characterization of b0 precipitate phase in Mg-2 at% Y alloy aged to peak hardness condition by high-angle annular detector dark-field scanning transmission electron microscopy (HAADFSTEM), Mater. Trans. 48 (2007) 84e87. [33] Y. Gao, H. Liu, R. Shi, N. Zhou, Z. Xu, Y.M. Zhu, J.F. Nie, Y. Wang, Simulation study of precipitation in an Mg-Y-Nd alloy, Acta Mater. 60 (2012) 4819e4832. [34] Z. Xu, M. Weyland, J.F. Nie, On the strain accommodation of b1 precipitates in magnesium alloy WE54, Acta Mater. 75 (2014) 122e133. [35] H. Liu, Y. Gao, Y.M. Zhu, Y. Wang, J.F. Nie, A simulation study of b1 precipitation on dislocations in an Mg-rare earth alloy, Acta Mater. 77 (2014) 133e150. [36] Z. Xu, M. Weyland, J.F. Nie, Shear transformation of coupled b1/b0 precipitates in Mg-RE alloys: a quantitative study by aberration corrected STEM, Acta Mater. 81 (2014) 58e70. [37] E.J. Kirkland, R.F. Loane, J. Silcox, Simulation of annular dark field stem images using a modified multislice method, Ultramicroscopy 23 (1987) 77e96. [38] K.H.W. Bos, A.D. Backer, G.T. Martinez, N. Winckelmans, S. Bals, P.D. Nellist, S.V. Aert, Unscrambling mixed elements using high angle annular dark field scanning transmission electron microscopy, Phys. Rev. Lett. 116 (2016) 246101. [39] J. Tan, Y.H. Sun, H.B. Xie, B.Z. Sun, Y. Qi, Atomic-resolution investigation of Yrich solid solution with an invariable orientation in Mg-Y binary alloy, J. Alloy. Comp. 766 (2018) 716e720. , R. Ferro, Phase equilibria in the binary rare[40] A. Saccone, S. Delfino, D. Maccio earth alloys: the erbium-magnesium system, Mater. Trans. A 23 (1992) 1005e1012. [41] S. Gorsse, C.R. Hutchinson, B. Chevalier, J.F. Nie, A thermodynamic assessment of the Mg-Nd binary system using random solution and associate models for the liquid phase, J. Alloy. Comp. 392 (2005) 253e262. [42] J.X. Zheng, B. Chen, Atomic-scale characterization of the equilibrium b phase in Mg-Nd-Y alloy by means of HAADF-STEM, Scanning 38 (2016) 743e746.
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Re-recognition of the aging precipitation behavior in the MgeSm binary alloy Hongbo Xie a, Boshu Liu a, Junyuan Bai a, Changli Guan a, Dongfang Lou a, Xueyong Pang a, Hong Zhao a, Shanshan Li a, Yuping Ren a, b, *, Hucheng Pan a, **, Changlin Yang c, Gaowu Qin a, b a
Key Laboratory for Anisotropy and Texture of Materials (Ministry of Education), School of Materials Science and Engineering, Northeastern University, Shenyang, 110819, China Research Center for Metallic Wires, Northeastern University, Shenyang, 110819, China c State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi'an, 710072, China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 5 June 2019 Received in revised form 10 September 2019 Accepted 16 September 2019 Available online 17 September 2019
Generally, the MgeSm system is one of the magnesium light rare-earth alloys (Mg-LRE), and its aging precipitation behavior is similar to that of the MgeNd or MgeCe alloys. For a long time, the b phase, precipitated after the b1 intermediate phase, has been considered to be Mg12Sm, although this has never been experimentally verified. In this study, however, we clearly confirm that the b phase is Mg5Sm with a face-centered cubic structure (space group: F43m; a ¼ 22.34 Å) by aberration-corrected scanning transmission electron microscopy (STEM) combined with quantitative STEM simulation. This is similar to the case of magnesium heavy rare-earth alloys (Mg-HRE). The lenticular-shaped b plates have a {1010}a habit plane and is fully coherent with the a-Mg matrix, with the orientation relationships among the b phase, b1 phase, and a-Mg matrix being (111)b//(111)b1//(1120)a and [110]b//[110]b1//[0001]a. This finding improves our understanding of the precipitation/strengthening mechanism in the MgeSm based alloys and is expected to guide the future design of novel high-strength heat-resistant Mg-RE based alloys. © 2019 Elsevier B.V. All rights reserved.
Keywords: Magnesium alloys Aging Precipitation Crystal structure HAADF-STEM Simulation
1. Introduction Magnesium alloys have long been considered as one of the potential weight-saving light structural materials [1e12]. However, their relatively poor mechanical properties, low ductility, and poor heat-resistant behaviors restrict their widespread applications [1e5]. Recently, Mg rare-earth (RE) alloys have received considerable attention due to their high strength at either room or elevated temperatures, excellent creep properties, and good formability [9,10]. The improvement of their mechanical properties is mainly ascribed to their high-density precipitates on the non-basal planes. However, the precipitation behaviors of these alloys are complex and closely related to those of the RE types [9,13]; the structures of the intermediate precipitates change with time at high
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (Y. Ren), [email protected] (H. Pan). https://doi.org/10.1016/j.jallcom.2019.152320 0925-8388/© 2019 Elsevier B.V. All rights reserved.
temperatures, resulting in the deterioration of their mechanical properties during service. In general, Mg-RE alloys can be classified simply into two groups based on the difference in the atomic numbers and chemical properties of the RE elements [9,13]: (i) Mg-LRE alloys, including Mg-Ce [14,15], Mg-Nd [16e19], and MgeSm [20e23]. The precipitation sequence of the Mg-LRE alloys is commonly agreed to be in the order of supersaturated solid solution (S$S$S$S) / b’’ (D019, Mg3RE) / bS’ (bco, Mg7RE) / b1 (fcc, Mg3RE) / b (bct, Mg12RE) [9,13e21]. (ii) Mg-HRE alloys, including Mg-Gd [24e30], Mg-Dy [31], and MgeY [25,29,32]. The precipitation sequence of the MgHRE alloys is generally accepted to be in the order of S$S$S$S / b’’ (D019, Mg3RE) / bL’ (bco, Mg7RE) / b1 (fcc, Mg3RE) / b (fcc, Mg5RE/cubic, Mg41RE5) [9,24e32]. From the precipitation sequences, it can be seen that one of the noticeable differences between the Mg-LRE and Mg-HRE alloys is the structure of the b0 intermediate phase, which is the main strengthening precipitate phase on the {1010}a plane. For the MgLRE alloys, the b0 precipitate is determined to have the short-
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structure b’ (denoted as bS’), having a composition of Mg7RE and an ordered orthorhombic structure (a ¼ 6.42 Å, b ¼ 11.12 Å, c ¼ 5.21 Å) [9,16e23]. For the Mg-HRE alloys, in contrast, the b0 phase has the long-structure (denoted as bL’), also with a composition of Mg7RE but an ordered orthorhombic structure (a ¼ 6.42 Å, b ¼ 22.23 Å, c ¼ 5.21 Å) [9,25e27,29e32], and the difference is the lattice parameter of b. However, notably, recent studies have indicated that both the bS’ and bL’ precipitates could co-exist in either MgeSm or MgeGd binary alloys [23,26,27]. Another difference in the aging precipitate sequence is the structure of the b precipitate after the b1 phase. For the Mg-HRE alloys, such as the MgeGd system, the b phase has been identified as the Mg5Gd phase with a face-centered cubic structure [9,24,28]. For the Mg-LRE alloys, however, all the phases precipitated after the b1 are thought to be the rod-shaped Mg12RE phase, having a body-centered tetragonal structure, such as Mg12Ce, Mg12Nd, and Mg12Sm [9,23]. Unfortunately, the b precipitate of Mg12Sm has not been confirmed by any transmission electron microscopy (TEM) or high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) studies to date. Therefore, this study aims to use aberration-corrected HAADFSTEM, selected area electron diffraction (SAED), and quantitative STEM simulation to characterize the b phase precipitated in a binary Mg-3wt.%Sm alloy during isothermal aging. Additionally, we unambiguously confirm that the b phase is a new Mg5Sm intermetallic compound (fcc, space group: F43m; a ¼ 22.34 Å), which is an unusual phenomenon in Mg-RE alloys that until now, had been thought unlikely. This finding improved our understanding of the precipitation/strengthening mechanism in MgeSm based alloys, and it could facilitate relevant modeling and simulation studies for this alloy system.
2. Experimental procedures The alloy with nominal composition of Mg-3wt.%Sm (Mg0.5 at.%Sm) was prepared by melting the pure Mg (99.9 wt%) and Mg-25Sm (wt.%) master alloy in the induction furnace with protection of the argon atmosphere. The molten alloy was stirred and kept at 760 C for 5 min and poured into a steel mold preheated to 300 C. The chemical composition of the obtained as-cast ingot was measured by OPTIMA 4300 DV composition analyzer, and the actual compositions was determined to be Mg-2.84Sm (wt.%). The
as-cast samples were solution treated at 520 C for 12 h, followed by water quenching and ageing in oil bath at 200 C for 300 h, and finally continued isothermally aged at 300 C for 48 h. The TEM specimens with a diameter of 3 mm were prepared by twin jet electro-polishing at - 40 C in mixture solution of 5.3 g lithium chloride, 11.2 g magnesium perchlorate, 500 ml methanol and 100 ml 2-butoxy ethanol, and subsequently ion milling with low energy electron beam. Finally, Gatan SOLARUS (950) Plasma Cleaning System was used to clean up the sample surfaces. TEM and HAADF-STEM observations were carried out using the JEMARM200F at an accelerating voltage of 200 kV, equipped with probe Cs corrector and cold field emission gun. The probe convergence is 25 mrad which yields a probe size of less than 0.1 nm, and the camera length was set to 8 cm which yields a collection semiangle of 48e327 mrad. The present study employed the quantitative TEM/STEM simulations software to simulate HAADF-STEM images, and the corresponding parameters used for simulations are identical with the experimental settings: a voltage of 200 kV, the third order spherical aberration C3 ¼ 0.05 mm, defocus values df ¼ 13.7 nm, beam convergence angle a ¼ 21.4 mrad.
3. Results and discussions Fig. 1a shows a low-magnification HAADF-STEM image recorded from the Mg-3Sm alloy isothermally aged at 200 C for 300 h, viewed along the [0001]a direction. It can be seen that there are high-density precipitates with high contrast in the a-Mg matrix, and these lenticular-shaped precipitates with a length more than ~200 nm distribute uniformly in three directions: ½1120a, ½1210a, and ½2110a. Furthermore, the majority of the precipitates in this aging-stage were identified as the b1 phase (Mg3Sm, fcc, space group: Fm3m; a ¼ 0.74 nm), which was reported in many Mg-RE based alloys [9,33e36]. Noteworthily, the orientation relationships between the b1 phase and the a-Mg matrix are [110]b1// [0001]a; (111)b1//(1120)a [9,33,35]. The atomic-scale HAADF-STEM image of one precipitate in Fig. 1a is enlarged and shown in Fig. 1b, and the b-Mg5Sm phase clearly forms in the b1 intermediate precipitate (as shown by the cyan circle). This observation result is consistent with the results of a recent study in the binary MgeGd and the WE54 alloys [28,34], a similar orientation relationships exists between the b-Mg5RE phase and the a-Mg matrix, i.e., [110]b//[0001]a; (111)b//(1120)a [9,24,28,34]. Additionally, the
Fig. 1. HAADF-STEM images of the Mg-3Sm alloy isothermally aged at 200 C for 300 h. (a) Low-magnification HAADF-STEM image. (b) A bright-contrast precipitate in Fig. 1a is enlarged and shown in Fig. 1b; the atomic-scale HAADF-STEM image indicated that there are some b-Mg5Sm phases formed in the over-aged state. The electron beam is parallel to the [0001]a direction.
H. Xie et al. / Journal of Alloys and Compounds 814 (2020) 152320
lattice image of the b phase is marked with a yellow rhombus in Fig. 1b, and the inset is its corresponding modeled atomic arrangement. At present, it is generally accepted that the b1 phase forms prior to the b phase, as evidenced by literature [9,24,28,34]. Our experimental observation indicates that the b phase nucleates in the existing b1 precipitate and grows gradually by the inside-out mode. It is generally believed that the Mg5RE phase can only precipitate in Mg-HRE alloys (MgeGd) during isothermal aging, and the terminal phase of the aging precipitation in binary MgeSm alloys is considered to be the Mg12Sm phase, instead of Mg5Sm. In order to characterize the detailed structure of the Mg5Sm phase observed above, and to confirm its thermal stability and evolution, the agedsample above was further isothermally aged at 300 C for 48 h. The TEM bright-field images and SAED patterns of the Mg-3Sm alloy further isothermally aged at 300 C for 48 h, viewed along the [0001]a, and [1120]a zone axes, respectively, are shown in Fig. 2. The lenticular-shaped precipitates with low contrast are observed to distribute uniformly in the ½1120a, ½1210a, and ½2110a directions, and coarsen up to a length of more than 500 nm and a thickness of ~100 nm. The corresponding SAED patterns, as shown in Fig. 2c and d, exhibit extra diffraction spots (labeled by yellow numbers) between the a-Mg matrix diffraction spots (labeled by red numbers). This indicates that the precipitates have an fcc structure, and it can be identified as the Mg5Sm phase (fcc, space group: F43m; a ¼ 22.34 Å), which is consistent with the Mg5Gd
3
equilibrium phase [28,34]. Besides, from the superimposed precipitate and matrix SAED patterns it can be determined definitely that the orientation relationships between the b phase and the aMg matrix are (111)b//(1120)a and [110]b//[0001]a. Furthermore, this implies that the orientation relationships among the b phase, b1 phase, and a-Mg matrix are (111)b//(111)b1//(1120)a and [110]b// [110]b1//[0001]a. Additional details concerning the b phase were obtained by conducting the HAADF-STEM analysis. Fig. 3a provides a lowmagnification HAADF-STEM image of the Mg-3Sm alloy further isothermally aged at 300 C for 48 h, viewed along the [0001]a direction, showing the b phase with a {1010}a habit plane (high contrast) precipitated in the a-Mg matrix, which is consistent with the bright-field TEM image shown in Fig. 2a. Fig. 3b shows two b precipitates connected at an angle of 120 , and the connection area, marked by a cyan circle in Fig. 3b, is further enlarged and presented in Fig. 3c. It can be seen that the b/b interface is incoherent. Fig. 3d presents a continuous b precipitate containing two kinds of contrasts, and the connection area marked by the cyan circle in Fig. 3d is further enlarged and shown in Fig. 3e. The local part of a b precipitate is further enlarged, as shown in Fig. 3f, and the atomic-scale HAADF-STEM image shows that there are many bright dots inside the b precipitate. Each bright dot represents a column rich in Sm atoms because the brightness of individual atomic columns in the HAADF-STEM images approximates to the square of the average atomic numbers (the atomic numbers for Mg and Sm are 12 and 62,
Fig. 2. Bright-field TEM images and the corresponding SAED patterns of the Mg-3Sm alloy with continued isothermal aging at 300 C for 48 h. (a and b) Bright-field TEM images. (c and d) Corresponding SAED patterns, and the cyan circles in Fig. 2a and b correspond to the SAED areas. The electron beam is parallel to the [0001]a (a and c) and [1120]a (b and d) directions.
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Fig. 3. HAADF-STEM images of the Mg-3Sm alloy subjected to continuous isothermal aging at 300 C for 48 h. (a) Low-magnification HAADF-STEM image. (b and c) HAADF-STEM image in Fig. 3b show two b precipitates connected an angle of 120 , and the connection area marked by a cyan circle is further enlarged and presented in Fig. 3c. (d and e) HAADFSTEM image in Fig. 3d shows a b precipitate containing two kinds of contrasts, and the connection area marked by a cyan circle is further enlarged and shown in Fig. 3e. (f) Atomicscale HAADF-STEM image. The electron beam is parallel to the [0001]a direction.
respectively) [37e39]. It can be seen that a periodic unit-cell with a 70.5 rhombus structure is marked by yellow dotted lines in this image, and the careful measurement of the HAADF-STEM image indicates that the value of the lattice parameter, a, is ~22.34 Å. The inset in Fig. 3f is the modeled atomic arrangement for the b precipitate viewed along the [110]b direction. In this model, Mg atoms are shown with red spheres and Sm atoms displayed with blue spheres, and it can be seen that the marked Sm atom arrangement shows perfect conformity with the HAADF-STEM results, both in symmetry and in interatomic distance. This verifies that the b phase precipitated in the binary MgeSm alloys is the Mg5Sm intermetallic compound. The Mg-3Sm aged alloy was further analyzed by HAADF-STEM viewed along the ½1120a direction, as shown in Fig. 4. Fig. 4a provides a low-magnification HAADF-STEM image, and the local part of this image, marked by a cyan square frame, is further enlarged and shown in the inset; the perfect match of the (110)b/ (0001)a planes indicates that the relationship of the b/a-Mg interface is fully coherent. The atomic-scale HAADF-STEM image is shown in Fig. 4b, in which a periodic lattice with an equilateral hexagon structure and a side length of ~9.12 Å is marked by yellow dotted lines; the inset in Fig. 4b shows that the ½111b modeled atomic arrangement of the b-Mg5Sm phase is in excellent agreement with the experimental observations above. Fig. 5 presents the HAADF-STEM simulation results for the Mg5Sm phase. The three-dimensional view of the unit cell is shown in Fig. 5a. It can be seen that the Mg5Sm phase has an fcc structure with a side length of 22.34 Å, and the corresponding simulated HAADF-STEM results viewed along the [110]b and ½111b directions are shown in Fig. 5b and c, respectively. The unique characteristics of the b phase, i.e., the rhombus lattice images (in the [110]b direction) and equilateral hexagon lattice images (in the ½111b direction) that consist of the bright dots, are clearly visible in the
simulated HAADF-STEM results, which is well consistent with the experimental HAADF-STEM observations shown in Figs. 3 and 4. This indicates that the b-Mg5Sm phase can be generated and stably exist. The Mg12RE phases, including the b-Mg12Nd, was originally regarded as the equilibrium phase but is, in fact, a metastable phase, and the equilibrium precipitate phase is be-Mg41Nd5 [9,40e42]. This is consistent with the MgeNd binary phase diagram. According to the MgeSm binary phase diagram, the equilibrium phase should be Mg41Sm5 (space group: I4/m; a ¼ 14.74 Å, c ¼ 10.40 Å); however, similar to the Mg12Sm phase, the Mg41Sm5 phase has never been confirmed by any electron diffraction or TEM techniques in the aging precipitation sequence. The age-hardenable MgeSm based alloy has relatively high mechanical properties attributed to its high-density b0 and b1 intermediate precipitates on the non-basal planes of the a-Mg matrix; however, its precipitation sequence is different from the previous one [20e23]. Since the precipitation behavior in the earlyor peak-aged stages is similar to those of the MgeCe and MgeNd alloys, its over-aged precipitation behavior was considered to be consistent with the Mg-LRE alloys, i.e., the Mg12Sm precipitate formed after the b1 phase. However, we unambiguously confirmed that the b phase precipitated in the over-aged MgeSm binary alloys is Mg5Sm, a new phase with a structure similar to that of Mg5Gd, not Mg12Sm. In other words, the over-aged precipitation behavior of the MgeSm alloys is consistent with that of the Mg-HRE alloys [9,24e30]. Additionally, the bS’ and bL’ intermediate precipitates have been confirmed to co-exist in the MgeSm alloys [23]. Therefore, from the viewpoint of the precipitation behavior, the MgeSm system should be classified as a Mg middle rare-earth alloy between the Mg-LRE and Mg-HRE alloys. Conversely, although the b’ and b1 intermediate phases have significant strengthening effects in the MgeSm alloys, their easy coarsening and change into the b
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Fig. 4. HAADF-STEM images of the Mg-3Sm alloy subjected to continuous isothermal aging at 300 C for 48 h. (a) Low-magnification HAADF-STEM image. (b) Atomic-scale HAADFSTEM image. The electron beam is parallel to the [1120]a direction.
precipitation in the MgeSm binary alloy, was investigated in this work. The b phase is the Mg5Sm intermetallic compound (fcc, space group: F43m; a ¼ 22.34 Å), which has a structure that is consistent with that of Mg5Gd, not that of the presently accepted Mg12RE. The lenticular-shaped b plates with a {1010}a habit plane distribute uniformly in three directions: ½1120a, ½1210a, and ½2110a, which is fully coherent with the a-Mg matrix. The orientation relationships among the b phase, b1 phase, and a-Mg matrix are (111)b//(111)b1// (1120)a and [110]b//[110]b1//[0001]a. In addition, the precipitation sequence of the MgeSm binary alloys should be modified as follows: S$S$S$S / intermediate phases / b1 (fcc, Mg3Sm) / b (fcc, Mg5Sm). The b1 phase has poor thermal stability at 300 C. This finding is important for understanding the precipitation mechanism of the MgeSm binary alloy, and is expected to guide the future design of novel high-strength MgeSm based alloys with high thermal stability. Acknowledgements The authors are grateful to the financial support from the National Key Research and Development Program of China (Grant No. 2016YFB0701202), the National Natural Science Foundation of China (Grant No. 51371046, No. 51525101, No. 51501032, No. U1610253), the Fundamental Research Funds of the Central Universities (Grant No. N141008001, No. N170204011), the Project of Promoting Talents in Liaoning Province (No. XLYC1808038), and the fund of the state Key Laboratory of Solidification Processing in NPU (Grant No. SKLSP201920). The experiments were conducted at the Institute of Advanced Material Technology of Northeastern University and we acknowledge the use of the aberration-corrected transmission electron microscopy. Fig. 5. HAADF-STEM simulation results of the Mg5Sm phase. (a) Three-dimensional view of the unit cell of the Mg5Sm phase. (b) Simulated [110]b HAADF-STEM image. (c) Simulated ½111b HAADF-STEM images.
phase would deteriorate the mechanical properties of the allow. Improving the thermal stability of these intermediate phases would be a challenging topic in the development of Mg-Sm-based alloys.
4. Conclusions The structure of the b phase, which is the terminal of the aging
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