Journal of Materials Science & Technology, 2021, 61(0): 114-118 DOI: 10.1016/j.jmst.2020.05.045

Research Article

Interfacial dislocations dominated lateral growth of long-period stacking ordered phase in Mg alloys

Qianqian Jina, Xiaohong Shao,a,*, Shijian Zhengc, Yangtao Zhoua, Bo Zhanga, Xiuliang Ma,a,b,**

aShenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang, 110016, China

bSchool of Materials Science and Engineering, Lanzhou University of Technology, Lanzhou, 730050, China

cSchool of Materials Science and Engineering, Research Institute for Energy Equipment Materials, Hebei University of Technology, Tianjin, 300130, China

Corresponding authors: * E-mail addresses:xhshao@imr.ac.cn(X. Shao),** Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang, 110016, China. E-mail addresses:xlma@imr.ac.cn(X. Ma).

Received: 2020-03-21   Accepted: 2020-05-17   Online: 2021-01-15

Abstract

Understanding the interface between strengthening precipitates and matrix in alloys, especially at the atomic level, is a critical issue for tailoring the precipitate strengthening to achieve desired mechanical properties. Using high-resolution scanning transmission electron microscopy, we here clarify the semi-coherent interfaces between the matrix and long-period stacking ordered (LPSO) phases, including 18R and 14H, in Mg-Zn-Y alloys. The LPSO/Mg interface features the unique configuration of the Shockley partial dislocations, which produces a near zero macroscopic strain because the net Burgers vectors equal zero. The 18R/Mg interface characterizes a dissociated structure that can be described as a narrow slab of 54R. There are two dislocation arrays accompanied to the 18R/54R and 54R/Mg interface, resulting a slight deviation (about 2.3°). The 14R/Mg interface exhibits the dislocation pairs associated with solute atoms. We further evaluate the stability and morphology of the corresponding interfaces based on elastic interaction, via calculating the mutual strong interactions between dislocation arrays, as well as that between the dislocations and solute atoms. The synchronized migration of interfacial dislocations and solute atoms, like move-drag behavior, dominates the lateral growth of LPSO phases in Mg alloys.

Keywords: Magnesium alloys ; Long-period stacking ordered phases ; Dislocations ; Heterogeneous interface

PDF (1588KB) Metadata Metrics Related articles Export EndNote| Ris| Bibtex  Favorite

Cite this article

Qianqian Jin, Xiaohong Shao, Shijian Zheng, Yangtao Zhou, Bo Zhang, Xiuliang Ma. Interfacial dislocations dominated lateral growth of long-period stacking ordered phase in Mg alloys. Journal of Materials Science & Technology[J], 2021, 61(0): 114-118 DOI:10.1016/j.jmst.2020.05.045

1. Introduction

Precipitate strengthening is one of the well-known strengthening mechanisms for materials, especially for magnesium alloys [[1], [2], [3], [4], [5], [6]]. This mechanism is based on controlling the distribution and the morphology of the precipitates in order to block the dislocation motion [1]. The heterogeneous interfaces between the precipitates and matrix are thus very important to determine the distribution and evolution of precipitates, consequently, which will strongly affect their strengthening of the materials. Understanding the structural features of these interfaces, particularly at the atomic level, will help us tune the microstructure of precipitates and meanwhile optimize the mechanical properties of the materials.

Magnesium alloys have attracted increasing attention due to their high specific strength and stiffness. Recently, long-period stacking ordered (LPSO) phases [[7], [8], [9], [10], [11], [12], [13], [14], [15]] were documented to play a significant role in strengthening magnesium alloys both at room and high temperatures [[16], [17], [18], [19]]. LPSO phases of micrometer-scale are frequently observed along grain boundaries, while those of nanometer-scale are distributed in the magnesium matrix [20]. They always have an orientation relationship with the α-Mg matrix: [11$\bar{2}$0]LPSO//[11$\bar{2}$0]matrix, (0001)LPSO//(0001)matrix. Regarding the interface between the LPSO phase and the α-Mg matrix, (0001)LPSO//(0001)matrix coherent interface has been demonstrated in detail [21,22], while the other semi-coherent (SC) interface perpendicular to the basal plane is less concerned. However, it would largely affect the lateral growth and morphology of the LPSO phase and in consequence, the mechanical performance of magnesium alloys.

In this work, we will explore the (01$\bar{1}$0) heterogeneous interface between the α-Mg matrix and commonly observed 18R and 14H LPSO phases in the Mg-Zn-Y alloy at the atomic scale. The mechanisms responsible for the featured interfaces are also proposed based on the cooperation of the strong interaction between ordered dislocations and the dislocations-solute atoms reactions.

2. Experimental

A near-equilibrium Mg97Zn1Y2 (at.%) alloy was fabricated by melting high purity Mg, Zn and Mg-30 wt% Y master ingots, and cooling down to room temperature at the rate of 4 K/min in a resistance furnace, under a protective mixed gas of SF6 (0.5 vol.%) and CO2 (99.5 vol.%). Thin foils for TEM observations cut from as-cast ingots were ground mechanically, electro-polished by the twin-jet polishing facility, and further thinned by the low energy ion beam. The high angle annular dark field-scanning transmission electron microscope (HAADF-STEM) imaging was performed by an aberration-corrected Titan3TM G2 60-300 scanning transmission electron microscopes equipped with a high-brightness field-emission gun (X-FEG) and double Cs correctors from CEOS, operated at 300 kV.

3. Results and discussion

3.1. Microstructure of interface between 18R and Mg matrix

18R LPSO phases are the dominant LPSO phases in the near-equilibrium Mg97Zn1Y2 (at.%) alloy. A low-magnification HAADF-STEM micrograph presents the step-like 18R/Mg interfaces, with (0001) coherent boundaries and (01$\bar{1}$0) SC boundaries, as shown in Fig. 1(a). 18R LPSO phase then exhibits bright contrast due to the enrichment of Zn and Y atoms, while α-Mg matrix shows dark contrast [23]. The atomic arrangement and compositional contrast of this (01$\bar{1}$0) SC boundaries are clearly presented in Fig. 1(b). We refer to the specific structure between 18R and matrix as 54R according to our previously introduced methods [9]. Two arrays of partial dislocations were detected to locate on (0001) slip planes with a repeated sequence b1, b2, and b3 denoted in Fig. 1(b). The 54R/Mg interface is made up of pure edge (90°) Shockley partial dislocations, whose Burgers vector b1 = 1/3[0$\bar{1}$10]. By contrast, the 18R/54R interface consists of b2 and b3 mixed (30°) Shockley partial dislocations. No tilt or twist misorientation between 18R and Mg matrix was detected based on Kikuchi line patterns, which indicates that the sum of both edge components and screw components of b1, b2 and b3 is zero. Thus, we deduced that b2 = 1/3[10$\bar{1}$0] and b3 = 1/3[$\bar{1}$100] (or b2 = 1/3[$\bar{1}$100] and b3 = 1/3[10$\bar{1}$0]) with their screw components of opposite sign. These dislocations produce a near zero macroscopic strain. Furthermore, there’s a small tilt angle, approximately 2.3°, between (0001)54R and (0001)Mg, or between (0001)54R and (0001)18R. Such a tilt angle results from two dislocation arrays in the 54R/Mg or 54R/18R interfaces, respectively. It can be estimated by Frank’s formula [24]: $\theta ={{b}_{\text{e}}}/D$, where ${{b}_{\text{e}}}={{b}_{1\text{e}}}={{b}_{2\text{e}}}+{{b}_{3\text{e}}}=\frac{1}{3}\left| \left[ 01\bar{1}0 \right] \right|=0.185\text{nm}$ is the magnitude of the edge component of dislocations in one period, and $D={{c}_{18\text{R}}}=4.69\text{nm}$ is the magnitude of one period of 18R LPSO phase. The atomic structural models of b1, b2 and b3 dislocations were proposed and atomic stacking sequences of Mg, 54R and 18R structures were illustrated (Fig. S1 in Supplementary Information), to interpret the role of each dislocation on the motion of atoms and the growth mechanism of 18R. The motion of b1, b2 and b3 dislocations leads to the stacking sequences transformation of Mg into 54R and 18R during the growth of 18R structure.

Fig. 1.

Fig. 1.   (a) Low-magnification HAADF-STEM image showing the 18R/Mg and 14 H/Mg coherent and semi-coherent (SC) interface. (b) Atomic-resolution HAADF-STEM image showing the complex 18R/Mg interfaces containing 54R in the near-equilibrium Mg97Zn1Y2 alloy. b1e, b2e and b3e are defined to be the edge component of b1, b2 and b3 dislocations, which are illustrated by open rectangles and white arrows. Two dislocation arrays in the 54R/Mg or 54R/18R interfaces result in the small tilt degree along the interfaces. The images were recorded along [2$\bar{1}\bar{1}$0]α zone axis.


3.2. Elastic interaction between dislocations at 18R/Mg interface

In order to estimate the interfacial stability of 18R and Mg matrix and to interpret the morphology of the complex boundaries, we calculated the elastic forces between b1, b2, and b3 dislocations based on elastic interaction. Generally, for two parallel bp and bq dislocations shown in Fig. 2(a), the interaction force Fpq is the sum of the force fpq,e between their edge components and the force fpq,s between their screw components [24]:

${{F}_{\text{pq}}}={{f}_{\text{pq},\text{e}}}+{{f}_{\text{pq},\text{s}}}$

Fig. 2.

Fig. 2.   (a) Forces on dislocation bq and solute atom due to stress field of dislocation bp. (b) Possible arrangement of b1, b2 and b3 dislocations responsible for transformation from 18R/Mg SC interface into 54R. (c) Glide force per unit length between neighboring dislocations, such as ${{F}_{11}}\left( x,3{{y}_{0}} \right)$, ${{F}_{22}}\left( x,{{y}_{0}} \right)$, ${{F}_{23}}\left( x,{{y}_{0}} \right)$, and their edge or screw components, like ${{f}_{23,\text{e}}}\left( x,{{y}_{0}} \right)$ and ${{f}_{23,\text{s}}}\left( x,{{y}_{0}} \right)$. (d) Force considered for the b2 and b3 walls with unlimited length acting on b1 dislocation array with length of 3ny0 expressed as ${{F}_{\omega }}=2n\sum\limits_{k=-\infty}^{+\infty}\,{{f}_{12}}\left( \omega ,\left( 3k+1 \right){{y}_{0}} \right)$, and the roles of free energy on the equilibrium width ω contributed by the interaction force. The value of x and ω is expressed in units of y0, while the values of force and free energy are expressed in units of $G{{b}^{2}}/(2\pi {{y}_{0}})$ and $G{{b}^{2}}/(2\pi )$, respectively.


Their absolute values of the force per unit length acting on bq by the stress field of bp evaluated at position (x, y) are calculated (S2). The component of force Fpq,x along the slip direction, x, determines the behavior of the dislocations since Shockley partial dislocations in hcp Mg structure slip on (0001) basal plane. Thus, for the sake of simplicity, we refer to the forces Fpq, fpq,e and fpq,s in the present work as the component of force Fpq,x, fpq,e,x, and fpq,s,x along the slip direction, x. The arrangement of dislocation arrays in the 18R/Mg SC interface before and after forming 54R structures is shown in Fig. 2(b). Then, the interaction forces between two adjacent dislocations and those contributed by their edge or screw components are plotted against the distance of the neighboring dislocations along slip direction x, in Fig. 2(c), such as ${{F}_{11}}\left( x,3{{y}_{0}} \right)$, ${{F}_{23}}\left( x,{{y}_{0}} \right)$, ${{f}_{23,\text{e}}}(x,{{y}_{0}})$, ${{f}_{23,\text{s}}}(x,{{y}_{0}})$ and ${{F}_{22}}(x,{{y}_{0}})$. For all the interaction forces marked in Fig. 2(c), the force is repulsive if $x\text{ }\cdot\text{ }{{F}_{\text{pq}}}\left( x \right)>0$ and vise versa. We can find that the stable equilibrium should be at $x=\pm \infty $, and $x={{x}_{0}}$ when ${{F}_{\text{pq}}}\left( {{x}_{0}} \right)=0$ and $\frac{\text{d}{{F}_{\text{pq}}}}{\text{d}x}\left( {{x}_{0}} \right)<0$. It follows that the tilt boundary associated with b1 array would be the most stable when all the b1 dislocations lie on the (01$\bar{1}$0) plane, which leads to the formation of a low-angle tilt boundary in the 54R/Mg interface. For the b2 and b3 dislocations, the interaction force ${{f}_{23,\text{s}}}\left( x,{{y}_{0}} \right)$ is several times larger than ${{f}_{23,\text{e}}}\left( x,{{y}_{0}} \right)$ if $-{{y}_{0}}<x<{{y}_{0}}$. Hence, the strong interaction force ${{F}_{23}}\left( x,{{y}_{0}} \right)$ between two neighboring b2 and b3 dislocations with opposite screw components would always be attractive. Notably, b2 and b3 dislocations have to be with opposite screw components and arrange alternately, as far as the stability of the interface is concerned. Such double-core dislocation array, consisting of b2 and b3 dislocations, is most stable when they lie on the (01$\bar{1}$0) plane, forming a low-angle tilt boundary in the 54R/18R interface.

To explain the stable width of 18R/Mg interface, we will analyze the equilibrium width of the 54R structure in the following via thermodynamics consideration of the complex boundaries with unit thickness, a length of 3ny0 (c18R) and a width of ω. The 54R with a width of ω is assumed to be transformed from ω/3 18R and 2ω/3 Mg via the two interfaces departing from each other by a relative distance of ω, as is shown in Fig. 2(b). The interaction force ${{F}_{12}}(x,y)$ (S2) between b1 and b2 (or b3) dislocations in 18R/Mg interfaces is

${{F}_{12}}(x,y)=-\frac{G{{b}^{2}}}{2\pi {{y}_{0}}}\frac{3{{y}_{0}}x\left( {{x}^{2}}-{{y}^{2}} \right)}{4{{\left( {{x}^{2}}+{{y}^{2}} \right)}^{2}}}$

For the interface with unlimited length, the force of b2 and b3 dislocation arrays acting on b1 dislocation array with an interval of 18R (3ny0), is ${{F}_{\omega }}=2n\sum\limits_{k=-\infty}^{+\infty}\,{{F}_{12}}\left( \omega ,\left( 3k+1 \right){{y}_{0}} \right)$, as plotted in Fig. 2(d). Obviously, the grain boundary is at an unstable equilibrium state as $x=0$, while it is at stable equilibrium as $x={{\omega }_{0}}=1.8{{y}_{0}}$ for 54R, merely considering the contribution of elastic force. The work done by the repulsion of the 18R/54R and 54R/Mg interfaces with a length of 3ny0 from 0 to ω, will decrease the free energy of the boundary. Further, the easier motion of b1 dislocations along + x zone due to the stronger interaction between solute atoms and b1 dislocations will slightly increase the width of 54R. This can rationalize the experimental results, that the width of observed 54R (3.4-7.8 nm, 2.2y0-5y0) is larger than the thermodynamic equilibrium width (1.8y0). The mutual reaction between solute atoms with dislocations was calculated in S3 in Supplementary Information. And the stronger interaction between solute atoms and b1 dislocation could be expected due to b1e = 2b2e = 2b3e, which triggers the easier motion of b1 dislocations and then the double-core dislocation b2 and b3.

3.3. Microstructure of interface between 14H and Mg matrix

Irregular-shaped 14H structures were also commonly observed in the near-equilibrium Mg97Zn1Y2 (at.%) alloy. The low-magnification HAADF-STEM image in Fig. 1(a) displays that 14 H/Mg SC boundaries do not necessarily parallel to (01$\bar{1}$0) plane, which was obtained with the incident electron beam parallel to [2$\bar{1}\bar{1}$0]Mg. To identify the detailed structure of the 14 H/Mg (01$\bar{1}$0) SC interfaces, an atomic-resolution HAADF-STEM image viewed along [2$\bar{1}\bar{1}$0]14H direction was shown in Fig. 3(a). AB'C'A and AC'B'A building blocks enriched in Zn/Y elements arrange alternately with a spacing of ${{y}_{0}}=7\times {{d}_{0002}}=1.82\text{nm}$. The two building blocks are associated with two 30° mixed Shockley partial dislocations b4 and b5, with edge components of Burgers vectors of same magnitude (${{b}_{4e}}={{b}_{5e}}=\frac{1}{6}\left| \left[ 01\bar{1}0 \right] \right|=0.092\text{nm}$) and opposite sign. Thus, the sum of the edge component of b4 and b5 is zero, i.e., no tilt angle between (0001)Mg and (0001)14H. The magnitude of the screw component of b4 and b5 is ${{b}_{4\text{s}}}=\frac{\sqrt{3}}{6}\left| \left[ 01\bar{1}0 \right] \right|=0.160\text{nm}$, and their signs of screw components can be determined to be of opposite sign, as there is no twist angle between (0001)Mg and (0001)14H. It means that the dislocations along 14 H/Mg also form a near zero macroscopic strain, like those along 18R/Mg. Close inspection of both images in Fig. 3(a) and (c) reveals that the tensile stress regions show brighter contrast, forming an inverted U, than the compressive stress regions along the 14 H/Mg interface. It strongly indicates that more Zn/Y atoms segregate in the tensile stress region. The intensity profile in Fig. 3(b) was obtained by integrating the image intensity in the rectangle in Fig. 3(a), which further clarifies that the intensity profile of the tensile stress region is higher than that of the compressive stress region. It should be mentioned that both flat interface in Fig. 3(a) and jogged interface in Fig. 3(c) were observed via atomic-resolution HAADF-STEM. The 14 H/Mg interfacial morphology is not restricted by the interfacial dislocation pairs, which should be closely related to the reaction between the dislocations and segregated solute atoms.

Fig. 3.

Fig. 3.   Atomic-resolution HAADF-STEM images viewed along [2$\bar{1}\bar{1}$0]α zone axis, showing the 14 H/Mg interfaces (a) approximately parallel to (01$\bar{1}$0) plane and (c) not parallel to (01$\bar{1}$0) plane. (b) Intensity profiles of the HAADF image in rectangle in (a).


3.4. Elastic interaction between dislocations and with segregated atoms at 14 H/Mg interface

To estimate its stability and to interpret its morphology, we also calculate the elastic forces between neighbor dislocations in 14 H/Mg interface. The dislocation configuration and formula are illustrated in Fig. S4 and S5 in Supplement Information. Fig. 4 presents the interaction forces ${{F}_{45}}\left( x,{{y}_{0}} \right)$ between two neighbor dislocations with screw components of opposite sign, b4 and b5. One can find that ${{F}_{45}}(x,{{y}_{0}})$ is always attractive, which makes the 14 H/Mg interface stable, as shown in the configuration of Fig. S4(b). Since Fig. 3 strongly suggested that solute atoms segregate to the tensile stress region of the dislocation array along 14 H/Mg interface, we calculated the elastic forces ${{F}_{I,x}}$ between a dislocation and solute atoms in its tensile stress region (Fig. S4), as plotted in Fig. 4. ${{F}_{I,x}}$ is obviously higher than F45. Hence, we presume that the dislocation pairs with the tensile stress region are restricted by the solute atoms and move synchronously. Namely, it is the interaction between solute atoms and dislocation pairs that dominates the lateral growth of 14 H/Mg interface.

Fig. 4.

Fig. 4.   Glide force per unit length acting along x zone between neighboring dislocations in 14 H/Mg interface and between their edge or screw components, such as ${{F}_{45}}\left( x,{{y}_{0}} \right)$, ${{f}_{45,\text{e}}}\left( x,{{y}_{0}} \right)$ and ${{f}_{45,\text{s}}}\left( x,{{y}_{0}} \right)$. Glide force per unit length ${{F}_{I,x}}$ acting along x zone between b4 dislocations and solute Zn/Y atoms in the tensile stress region, which can cause volume increase $\text{ }\Delta\text{ }V={{b}_{\text{e}}}{{y}_{0}}l$.


3.5. Comparison of migration mechanism of the LPSO/Mg interface and incoherent twin boundary

Our atomic scale images and elastic interaction suggest that the 18R/Mg and 14 H/Mg SC interfaces are dominated by the interfacial dislocations and dislocation-solute reactions. They are in accordance with the mechanism proposed for the interface-controlled mobility when the abundant solute reacts with the interface defects, such as dislocations and steps, etc. Additionally, the propagation mechanism for the LPSO/Mg interface exhibit analogies to the “move-drag” deformation twinning mechanism, which recently is proposed to explain the movement of the incoherent twin boundary (ITBs) in face-centered-cubic (fcc) metals [[25], [26], [27]]. Namely, one or two dislocations are triggered, and then the rest dislocations are dragged along, inducing the interface’s migration. The slight difference lies in that the driving force. The applied shear stress is for the ITBs, while the formation enthalpy is for the LPSO/Mg SC interface. Specifically, contrary to the shear strain-induced “move-drag” mechanism of ITBs, the interaction between dislocation arrays and solute atoms is assumed to dominate the LPSO/Mg SC interface in magnesium alloys. In detail, the unique dislocation configuration at the LPSO/Mg interface, as shown in Fig. 1, Fig. 3, lowers the interfacial energy. The stronger interaction between b1 dislocation arrays and solute atoms moves b1 dislocation arrays firstly, and then drag double-core dislocation (b2and b3) arrays, which is responsible for 18R/LPSO SC interface featuring (01$\bar{1}$0)18R//(01$\bar{1}$0)Mg; the stronger reaction between dislocation pairs and solute atoms in their mutual tensile stress region prompts the dislocation pairs to extend synchronously, and hence results in the wavy 14 H/LPSO boundary vertical to (0001) plane.

4. Conclusion

In summary, we unravel the morphology, dislocation configuration, and migration mechanism of 18R/Mg and 14 H/Mg SC boundaries in a near-equilibrium Mg97Zn1Y2 (at.%) alloys via a combination of aberration-corrected transmission electron microscopy and elastic mechanics calculation. The dislocations at the respective LPSO/Mg interface produce a near zero macroscopic strain since the net Burgers vectors is zero. The lateral growth of the LPSO/Mg interface is dominated by the unique interfacial dislocations, as well as the interaction between dislocations and solute atoms. The 18R/Mg SC interface is parallel to (01$\bar{1}$0) plane and dissociated into 54R structure, which is sandwiched between the pure edge dislocation (b1) array and the double-core dislocation (b2 + b3) array. The stronger interaction between solute atoms and b1 dislocation than b2/b3 leads to an easier motion of b1 dislocation array and then dragging b2/b3 dislocations. The irregular 14 H/Mg SC interface features an array of the dislocation pairs, which consist of two dislocations with equal but opposite Burgers vectors. These dislocations move synchronously due to the strong interaction between the dislocation pairs and solution atoms in their tensile stress region.

Acknowledgments

This work is supported financially by the National Natural Science Foundation of China (Nos. 51801214 and 51871222), and the Liaoning Provincial Natural Science Foundation (No. 2019-MS-335). The authors express their sincere gratitude to Prof. J. Wang, University of Nebraska-Lincoln, for valuable discussion.

Appendix A. Supplementary data

Supplementary material related to this article can be found, in the online version, at doi:https://doi.org/10.1016/j.jmst.2020.05.045.

Reference

J.F. Nie, Metall. Mater. Trans. A 43 (2012) 3891-3939.

DOI      URL     [Cited within: 2]

Y. Li, C. Yang, X. Zeng, P. Jin, D. Qiu, W. Ding, Mater. Charact. 141 (2018) 286-295.

DOI      URL     [Cited within: 1]

H. Liu, H. Huang, J.P. Sun, C. Wang, J. Bai, A.B. Ma, X.H. Chen, Acta Metall. Sin. (Engl. Lett.) 32 (2018) 269-285.

DOI      URL     [Cited within: 1]

X. Wei, L. Jin, F. Wang, J. Li, N. Ye, Z. Zhang, J. Dong, J. Mater. Sci. Technol. 44 (2020) 19-23.

DOI      URL     [Cited within: 1]

C. Zhu, B. Chen, Adv. Eng. Mater. 21 (2018), 1800734.

DOI      URL     [Cited within: 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, X.H. Chen, G. Chen, K.K. Deng, H.J.Y. Wang, J. Mater. Sci. Technol. 34 (2018) 245-247.

DOI      URL     [Cited within: 1]

Y.M. Zhu, A.J. Morton, J.F. Nie, Acta Mater. 58 (2010) 2936-2947.

DOI      URL     [Cited within: 1]

AbstractThe 18R and 14H long-period stacking ordered structures formed in Mg–Y–Zn alloys are examined systematically using electron diffraction and high-angle annular dark-field scanning transmission electron microscopy. In contrast to that reported in previous studies, the 18R structure is demonstrated to have an ordered base-centred monoclinic lattice, with Y and Zn atoms having an ordered arrangement in the closely packed planes. Furthermore, the composition of 18R is suggested to be Mg10Y1Zn1, instead of the Mg12Y1Zn1 composition that is commonly accepted. The 14H structure is also ordered. It has a hexagonal unit cell; the ordered distribution of Y and Zn atoms in the unit cell is similar to that in the 18R and its composition is Mg12Y1Zn1. The 18R unit cell has three ABCA-type building blocks arranged in the same shear direction, while the 14H unit cell has two ABCA-type building blocks arranged in opposite shear directions.]]>

E. Abe, A. Ono, T. Itoi, M. Yamasaki, Y. Kawamura, Philos. Mag. Lett. 91 (2011) 690-696.

DOI      URL     [Cited within: 1]

Q.Q. Jin, X.H. Shao, X.B. Hu, Z.Z. Peng, X.L. Ma, Philos. Mag. 97 (2017) 1-16.

DOI      URL     [Cited within: 2]

Q.Q. Jin, X.H. Shao, X.B. Hu, Z.Z. Peng, X.L. Ma, Philos. Mag. Lett. 97 (2017) 180-187.

DOI      URL     [Cited within: 1]

H. Zhang, C.Q. Liu, Y.M. Zhu, H.W. Chen, L. Bourgeois, J.F. Nie, Acta Mater. 152 (2018) 96-106.

DOI      URL     [Cited within: 1]

S.B. Mi, Q.Q. Jin, Scr. Mater. 68 (2013) 635-638.

DOI      URL     [Cited within: 1]

C. Liu, Y. Zhu, Q. Luo, B. Liu, Q. Gu, Q. Li, J. Mater. Sci. Technol. 34 (2018) 2235-2239.

DOI      URL     [Cited within: 1]

D. Egusa, E. Abe, Acta Mater. 60 (2012) 166-178.

DOI      URL     [Cited within: 1]

We propose structural models of the unique long period stacking/order (LPSO) phases formed in Mg-Zn-RE alloys, based on Z-contrast scanning transmission electron microscopy observations and first principles calculations. The LPSO structures are long period stacking derivatives of the hcp Mg structure, and the Zn/RE distributions are restricted at the four close-packed atomic layers forming local fcc stacking (i.e. a local ABCA stacking). Chemical order is well developed for the LPSO phases formed in Mg(97)Zn(1)Er(2) (14H type) and Mg(85)Zn(6)Y(9) (18R type) alloys with pronounced superlattice reflections, and the relevant Zn/RE distributions clearly emerge in the Z-contrast atomic images. Initial ternary ordered models were constructed by placing all the atoms at the ideal honeycomb sites, leading to plausible space groups of P6(3)/mcm for the 14H type and C2/m, P3(1)12 or P3(2)12 for the 18R type. The characteristic ordered features are well represented by local Zn(6)RE(8) clusters, which are embedded in the fcc stacking layers in accordance with the L1(2) type short-range order. Energy favored structural relaxations of the initial model cause significant displacement of the Zn/RE positions, implying that strong Zn-RE interactions may play a critical role in phase stability. The LPSO phases seem to tolerate a considerable degree of disorder at the Zn and RE sites with statistical co-occupations by Mg, extending the non-stoichiometric phase region bounded along the Zn/RE equiatomic line from similar to Mg(94.0)Zn(2.0)Y(4.0) to similar to Mg(83.3)Zn(8.3)Y(8.3). (C) 2011 Acta Materialia Inc. Published by Elsevier Ltd.

M. Yamasaki, M. Matsushita, K. Hagihara, H. Izuno, E. Abe, Y. Kawamura, Scr. Mater. 78-79 (2014) 13-16.

DOI      URL     [Cited within: 1]

A. Inoue, Y. Kawamura, M. Matsushita, K. Hayashi, J. Koike, J. Mater. Res. 16 (2001) 1894-1900.

DOI      URL     [Cited within: 1]

Y. Kawamura, K. Hayashi, A. Inoue, T. Masumoto, Mater. Trans. 42 (2001) 1171-1174.

[Cited within: 1]

X.H. Shao, Z.Q. Yang, X.L. Ma, Acta Mater. 58 (2010) 4760-4771.

DOI      URL     [Cited within: 1]

AbstractThe deformation behavior and corresponding microstructure evolution of a Mg97Zn1Y2 (at.%) alloy with a long period stacking ordered (LPSO) structure subjected to hot compression were investigated. The peak stress at 573 K was about 190 MPa, and no macroscopic fracture took place up to a strain of about 60%. The mechanisms responsible for the mechanical performance of the Mg97Zn1Y2 (at.%) alloy are discussed based on microstructural investigations using various electron microscopy techniques. The high strength at elevated temperature could be attributed to synergetic strengthening refinement of the LPSO via kinking and a limited fraction of dynamical recrystallization. Microcracks nucleated at the interfaces in the sandwich structure composed of LPSO and nanometer thick Mg slices could weaken the alloy at late stages of deformation, but their propagation could be limited within the individual kink band where the microcracks nucleated, which could ensure the capability of the alloy to resist premature or catastrophic fracture. Furthermore, lack of deformation twins in Mg grains effectively reduced the potential nucleation sites for cracks, which should be another reason for the good ductility of the alloy. These findings may provide or evoke insights into methods for optimizing the mechanical properties of Mg alloys.]]>

X.H. Shao, Z.Z. Peng, Q.Q. Jin, X.L. Ma, Acta Mater. 118 (2016) 177-186.

DOI      URL     [Cited within: 1]

Y. Kawamura, M. Yamasaki, Mater. Trans. 48 (2007) 2986-2992.

DOI      URL     [Cited within: 1]

Y.M. Zhu, A.J. Morton, J.F. Nie, Acta Mater. 60 (2012) 6562-6572.

DOI      URL     [Cited within: 1]

(alpha), 18R transforms in-situ to 14H during prolonged heat treatment at 500 degrees C. The 18R to 14H transformation is shown to occur most readily in regions where the 18R structure has irregularities in the building block stacking, in particular where a pair of adjacent building blocks is separated by four rather than two alpha-Mg layers. It is proposed that the diffusional-displacive 18R to 14H transformation rate is controlled by the diffusion rate of Y and Zn atoms into the segregation layers. (C) 2012 Acta Materialia Inc. Published by Elsevier Ltd.]]>

J.F. Nie, Y. Zhu, A.J. Morton, Metall. Mater. Trans. A 45 (2014) 3338-3348.

DOI      URL     [Cited within: 1]

Y.M. Zhu, M. Weyland, A.J. Morton, K. Oh-ishi, K. Hono, J.F. Nie, Scr. Mater. 60 (2009) 980-983.

DOI      URL     [Cited within: 1]

D. Hull, D.J. Bacon, Introduction to Dislocations, fifth ed., Butterworth-Heinemann, Boston, 2011, pp. 75-79.

[Cited within: 2]

J. Wang, O. Anderoglu, J. Hirth, A. Misra, X. Zhang, Appl. Phys. Lett. 95 (2009), 021908.

DOI      URL     [Cited within: 1]

J. Wang, A. Misra, J. Hirth, Phys. Rev. B 83 (2011), 064106.

DOI      URL     [Cited within: 1]

D. Hofmann, F. Ernst, Interface Sci. 2 (1995) 201-210.

[Cited within: 1]

/