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Phase coexistence in Ti6Sn5 intermetallics
Intermetallics 51 (2014) 48e52
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Intermetallics journal homepage: www.elsevier.com/locate/intermet
Phase coexistence in Ti6Sn5 intermetallics A.A. Oni, D. Hook, J.P. Maria, J.M. LeBeau* Department of Materials Science & Engineering, North Carolina State University, Raleigh, NC 27695-7907, USA
a r t i c l e i n f o
a b s t r a c t
Article history: Received 23 December 2013 Received in revised form 7 March 2014 Accepted 10 March 2014 Available online
Here we report the structural characterization of a complex TieSn intermetallic compound, Ti6Sn5. From X-ray diffraction, the resulting compound was observed to exist in both orthorhombic and hexagonal phases. Analysis by electron microscopy revealed that “planar-like” defects form throughout the material. Atomic resolution aberration-corrected scanning transmission electron microscopy reveals that these “planar-like” defects represent the coexistence of the orthorhombic and hexagonal phases within single grains. The resulting interwoven phases range in thickness from a fraction to multiple unit cells and exhibit coherent phase boundaries with the matrix grain. Ó 2014 Elsevier Ltd. All rights reserved.
Keywords: A. Intermetallics, miscellaneous D. Phase interfaces D. Defects: planar faults F. Electron microscopy, transmission F. Diffraction
1. Introduction Titanium alloys and intermetallics are utilized in many applications spanning aircraft production, medical instruments [1] and mechanical dampening [2]. They exhibit outstanding oxidation and corrosion resistance, low density, and mechanical properties [3,4]. Further, the electrical and magnetic properties of the intermetallics offer the potential for additional functionality. Experimental studies of hexagonal Ti6Sn5, for example, revealed a ground state in close proximity to a non-magnetic-magnetic phase boundary [5]. Moreover, doping Ti6Sn5 with rare earth elements introduces ferromagnetic instability and ordering [6]. Further, little is known about defects in Ti6Sn5 compounds, which may ultimately have important consequences on properties. The complex Ti6Sn5 intermetallic line compound has been reported to exhibit hexagonal, h-Ti6Sn5 (space group P63/mmc) [3], and orthorhombic, o-Ti6Sn5 (space group Immm) [7], phases. The lattice parameters for h-Ti6Sn5 and o-Ti6Sn5 are a ¼ 0.9248 nm, c ¼ 0.5690 nm [3] and a ¼ 1.693 nm, b ¼ 0.9144 nm and c ¼ 0.5735 nm [7], respectively. The phase diagram of the TieSn system [8] shows that hexagonal Ti6Sn5 (h-Ti6Sn5) is energetically favorable above 751 C, and the o-Ti6Sn5 phase is stable below [5]. Furthermore, the enthalpy of formation, DHf reported from firstprinciple calculations [9,10] for h-Ti6Sn5 and o-Ti6Sn5 is 37.25 kJ/mol and 34.92 kJ/mol, respectively [11]. Given the
* Corresponding author. Tel.: þ1 919 515 5049. E-mail address: [email protected] (J.M. LeBeau). http://dx.doi.org/10.1016/j.intermet.2014.03.002 0966-9795/Ó 2014 Elsevier Ltd. All rights reserved.
negligible difference in DHf, either phase may form depending on processing conditions. For example, experiments have demonstrated that the h-Ti6Sn5 phase is typically accompanied by orthorhombic o-Ti6Sn5 [5]. Defect analysis of complex intermetallic structures is challenging due to their large lattice parameters (small diffraction vector spacing) and because they possess a large number of atoms per unit cell, ranging from many tens of atoms up to more than a thousand [12,13]. The introduction of aberration correction in scanning transmission electron microscopy (STEM) has opened the doorway to the characterization of defects in these structures with unprecedented clarity [12,14,15]. In particular, high-angle annular dark-field (HAADF) STEM has proven an invaluable tool for materials characterization as its image intensities scale with atomic number. For example, aberration-corrected STEM was used to investigate stacking faults and dislocation cores in the Laves phase Cr2Hf, as a result, improving our general understanding of deformation mechanism and phase transformation kinetics in complex structures [14]. In this Article, we present a detailed structural analysis of the Ti6Sn5 intermetallic compound. By conventional BF/DF TEM, we show the formation of “planar-like” defects formed throughout grains of Ti6Sn5, which are reminiscent of stacking faults. We employ aberration-corrected high-angle annular dark-field scanning transmission electron microscopy (HAADF STEM) to determine the atomic scale structure of these defects. Structural models and HAADF-STEM image simulations are used to confirm that these “planar-like” defects result from the coexistence of h-Ti6Sn5 and oTi6Sn5 phases within single grains.
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2. Materials and methods Bulk TieSn mixtures were prepared from w20 mm Ti and w5 mm Sn powders combined in a 6:5 atomic ratio. The mixture was then homogenized with ball milling for 15 min. The powder was then compressed in a 1 cm die at 12 kpsi. The pellet was then heated to 1275 C for 17 h in flowing Ar ðpO2 z105 atmÞ. X-ray diffraction experiments were performed with a PANalytical Empyrean diffractometer equipped with a variable-slit scanning detector and Cu Ka (l ¼ 1.5418 nm) incident radiation. Specimens for electron microscopy were prepared by conventional mechanical polishing using an Allied High Tech MultiprepÔ. A Fischione Model 1050 Ar ion mill with a LN2 cooled stage was used to thin samples to electron transparency. Conventional bright-/dark-field images were acquired with a JEOL 2000FX transmission electron microscope operated at 200 kV and equipped with a LaB6 emitter. A probe-corrected FEI Titan G2 60-300 KV S/ TEM equipped with a high-brightness Schottky field emission gun was used for HAADF-STEM imaging at 200 kV. The convergence and collection semi-angles for the Titan S/TEM were 21 mrad and 76e 400 mrad respectively. Atomic resolution HAADF-STEM image simulations were performed using the multislice approach [16]. The images were simulated using the following input parameters: sample thickness ¼ 20 nm, C3 ¼ 6 mm, C5 ¼ 2.5 mm, and the same convergence and collection geometries that were used during experimental micrograph acquisition. Thermal diffuse scattering was introduced in the simulations using the frozen lattice model [17] and an estimated root-mean-squared (r.m.s.) displacement of 0.09 A for both Ti and Sn. Simulated images were convolved with a Gaussian function, 0.1 nm full-width half-max, to approximately account for the finite effective source size [18]. Simulated images of phase coexistence regions were generated by joining separate simulations of h-Ti6Sn5 and o-Ti6Sn5 for the same crystal thickness. 3. Results and discussion X-ray diffraction analysis of the as-processed bulk pellet, presented in Fig. 1, showed the formation of the Ti6Sn5 intermetallic with trace amounts of both Ti5Sn3 and Ti2Sn5. The peaks labeled ‘Other’ correspond to Ti5Sn3 and Ti2Sn5. The presence of the hTi6Sn5 and the o-Ti6Sn5 phases was confirmed by atomic resolution HAADF-STEM imaging as shown in Fig. 2. The h-Ti6Sn5 phase was the most frequently observed throughout the bulk of the specimen. The crystal structure of h-Ti6Sn5 consists of ABCB0 ABCB0 stacking sequence when viewed along [100] projection as seen in Fig. 2(a).
Fig. 1. Powder X-ray diffraction pattern of the TiSn alloy. h-Ti6Sn5 and o-Ti6Sn5 peaks are indicated, while traces of Ti2Sn5 and Ti5Sn3 phases detected are presented as Other phases.
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The A planes consist of atom columns containing both Sn atoms and Ti atoms. The B ‘planes’ exhibit alternating atom column doublets (TieSn) and triplets (TieSneSn). The C plane contains only Ti atom atoms. The B0 and B planes are related through a two-fold rotation about the C plane. The stacking sequence of the o-Ti6Sn5 phase along the [010] projection is nearly identical to h-Ti6Sn5 except that mirror symmetry is introduced across the A planes, such that the sequence becomes ABCB0 AB0 CB. Comparison of HAADF-STEM images from experiment and simulations along the hexagonal h100ih (Fig. 2(b)) shows the closely-spaced SneTi (doublet) columns have the brightest intensity and the pure Ti columns have the dimmest intensity. Furthermore, a similar intensity variation in atomic columns is noted in o-Ti6Sn5, in Fig. 2(c), however, there are important differences between the HAADF-STEM images of the two phases. A chevron-like pattern is observed for the hexagonal phase, as noted with white lines in Fig. 2(b), whereas mirror symmetry is present across the C planes of o-Ti6Sn5 (ABCB0 AB0 CB sequence). These image intensity details are essential for direct identification of the crystal structure and defect structures in this system. At the microstructural level, conventional transmission electron microscopy revealed planar defects throughout the Ti6Sn5 grains, as indicated in Fig. 3(a). Individual grains typically possessed a high density of these planar defects, with roughly 0.10 per nm. Furthermore, selected area diffraction patterns along [010] showed streaking of the diffraction spots parallel to g100, see Fig. 3(b). The observed streaks are indicative of randomly displaced (100) planar faults, and are typically observed with high stacking fault or twin boundary density structures [19]. Further, the planar-type defects were investigated with aberration-corrected STEM, where they appear as a disturbance in the stacking sequence and a corresponding decrease in the image intensity, as indicated by the white arrows in Fig. 3(c). Close examination of the atomic structure in Fig. 3(c) suggests the possibility that these defects are the formation of 180 domain boundaries, where the chevrons of the [100]h projections mirror across the defect plane relative to one another. Rather than abrupt boundaries, however, the defects exhibited nm-scale spatial extent perpendicular to the (100)h fault plane ranging from approximately 1e10 nm. Moreover, in regions where the defects appeared to extend beyond a single layer, atomic resolution HAADF imaging revealed the formation of multiple unit cells of o-Ti6Sn5 within an h-Ti6Sn5 matrix. An example of two unit cells of h-Ti6Sn5 embedded within an o-Ti6Sn5 is shown in the HAADF-STEM image in Fig. 3(d). These coexistence defects are consistent with the interpretation of o-Ti6Sn5 within an h-Ti6Sn5 matrix. From inspection of Fig. 3(c) and (d), image simulations qualitatively reproduce the features across the boundary. An averaged intensity line profile from experiment is presented in Fig. 4 and compared to the simulation. To increase the signal-to-noise ratio, the boxed region indicated in Fig. 3(c) was used as a template to find identical phase coexistence regions within the experimental image using a normalized crosscorrelation approach [20]. These regions were subsequently averaged and integrated along [001]h to generate the line profile. As shown, the image contrast is also in good qualitative agreement when employing the phase coexistence model. To complement the results from [100]h, the planar defects were also observed along [011]h, as seen in Fig. 5. The defects again extended beyond a single plane and the structure at the boundary was inconsistent with anti-phase or domain boundaries, indicating the need for an alternative structural model to explain the defect features. Rather, the introduction of a sub-unit layer or layers of oTi6Sn5 provides the necessary symmetry needed to account for the defect mirror plane and spatial extent. The phase coexistence model is shown schematically in Fig. 5. As indicated in the stacking
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Fig. 2. (a) HAADF-STEM image of h-Ti6Sn5 in the h100i projection. (b) HAADF-STEM image of o-Ti6Sn5 in the h010i projection. (c) Schematic showing stacking sequence of h-Ti6Sn5 and o-Ti6Sn5 structure. Atomic model and HAADF simulations are superimposed on the HAADF-STEM images.
sequence schematic, the stacking sequence for h-Ti6Sn5 in this projection is DED0 E0 D, where D0 and E0 are related to D and E through inversion symmetry across the sub-unit boundary. For comparison, the stacking sequence for o-Ti6Sn5 is simply DEDE. Along the [011] zone, the atom columns project a distinctly different atom positions along [011]h and [011]o. The formation of these defects in the Ti6Sn5 correlates with the minimal difference in DHf for h-Ti6Sn5 and o-Ti6Sn5, and are likely to result from processing conditions. As a result, these defects may not occur in all cases and will depend upon the kinetics in addition to the thermodynamic stability. Moreover, these results indicate the possibility that, rather than the formation of stacking faults or other planar defects, the coexistence of h-Ti6Sn5 and o-Ti6Sn5 acts
to relieve stress during mechanical deformation. Furthermore, given the close proximity of the hexagonal phase to non-magneticmagnetic phase boundary, these results may have important consequences on controlling the magnetic properties of Ti6Sn5 intermetallic compounds. Phase coexistence of hexagonal and orthorhombic phases has also been observed by varying composition [21], temperature [22], and pressure [23]. For example, structural defects arising from stacking variants have been observed in the Mg-based pseudo-binary Laves phases [24e26]. As we have observed for Ti6Sn5, Laves phases are capable of coexisting in the same system where they exhibit abrupt changes in the stacking sequence [27,28] and similar to Ti6Sn5, the Laves phases exhibit fundamental prototypes with a
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Fig. 3. (a) Two-beam dark-field TEM imaging of a Ti6Sn5 grain using g100 reflection. Planar defects are indicated with arrows. (b) [010] zone axis selected area diffraction pattern. The streaking of the diffraction spots are indicative of high stacking fault faults or twin boundary. (c) HAADF-STEM image of h-Ti6Sn5 in the h100i projection with o-Ti6Sn5 phase embedded within the matrix. (d) h-Ti6Sn5 phase embedded within o-Ti6Sn5.
range of crystal structures, namely cubic e C14, and hexagonal e C15 and C36 [27]. 4. Conclusions Through a combination of X-ray diffraction, conventional TEM, and aberration-corrected STEM, we have demonstrated that the identification of defects in complex intermetallic structures
Fig. 4. Average intensity line profile extracted from an HAADF-STEM image of an hTi6Sn5eo-Ti6Sn5 phase coexistence region (black, solid line e see indicated area in Fig. 3(c)) compared with simulation (red, dashed line). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 5. HAADF images of phase coexistence of h-Ti6Sn5 and o-Ti6Sn5 in the h011i projection. The inset is a schematic for the stacking sequence of h-Ti6Sn5 and o-Ti6Sn5 in the h011i projection.
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remains a challenging endeavor. Specifically, we have shown that conventional TEM analysis can lead to misidentification of planar defects when phase coexistence occurs. In the case of Ti6Sn5 planar defects, the coexistence of hexagonal and orthorhombic phases is readily revealed by atomic resolution imaging. Moreover, our results demonstrate the power of aberration-corrected STEM to fully characterize structural defects in complex structures. Additional work is required to relate the observed phase coexistence with the mechanical and magnetic properties of Ti6Sn5. Acknowledgments The authors acknowledge Oak Ridge National Laboratory ShaRE program for the preliminary data acquisition for this project. A.A.O and J.M.L acknowledge support from the Air Force Office of Scientific Research (Grant no. FA9550-12-1-0456). J.P.M and D.H. acknowledge funding for this project through DuPont Microcircuit Materials under contract 2008-2446. The authors acknowledge the use of the Analytical Instrumentation Facility (AIF) at North Carolina State University, which is supported by the State of North Carolina and the National Science Foundation. References [1] Wang X, Li W, Fang G, Wu C, Lin J. First-principles calculations on the electronic structure and cohesive properties of titanium stannides. Intermetallics 2009;17(9):768e73. [2] Wong CR, Fleischer RL. Low frequency damping and ultrasonic attenuation in Ti3Sn-based alloyssed alloys. J Mater Res 1994;9(6):1441e8. [3] Villars P, Calvert LD. Pearson’s handbook crystallographic data for intermetallic phases, vol. 2. Metals Park, OH: American Society for Metals; 1985. [4] Kim Y-W, Dimiduk D. Progress in the understanding of gamma titanium aluminides. JOM 1991;43(8):40e7. [5] Drymiotis F, Lashley JC, Fisk Z, Peterson E, Nakatsuji S. Physical properties of the b-Ti6Sn5 system. Philos Mag 2003;83(27):3169e78. [6] Jeong T. Electronic structure and magnetic properties of b-Ti6Sn5. J Magnet Magnet Mater 2007;309(1):71e4. [7] van Vucht JHN, Bruning HACM, Donkersloot HC, de Mesquita AHG. Philips Res Rep 19(5);407e21. [8] Okamoto H. SneTi (TineTitanium). J Phase Equilibria Diff 2010;31(2):202e3.
[9] Kresse G, Hafner J. Ab initio molecular-dynamics simulation of the liquidmetal-amorphous-semiconductor transition in germanium. Phys Rev B 1994;49:14251e69. [10] Kresse G, Furthmüller J. Efficient iterative schemes for ab initio total-energy calculations using a plane-wave basis set. Phys Rev B 1996;54:11169e86. [11] Colinet C, Tedenac J-C, Fries SG. Structural stability of intermetallic phases in the SneTi system. Calphad 2009;33(1):250e9. [12] Heggen M, Houben L, Feuerbacher M. Plastic-deformation mechanism in complex solids. Nat Mater 2010;9(4):332e6. [13] Urban K, Feuerbacher M. Structurally complex alloy phases. J Non-Crystall Solids 2004;334e335(15):143e50. [14] Chisholm MF, Kumar S, Hazzledine P. Dislocations in complex materials. Science 2005;307(5710):701e3. [15] Pyrz WD, Blom DA, Sadakane M, Kodato K, Ueda W, Vogt T, et al. Atomic-level imaging of Mo-V-O complex oxide phase intergrowth, grain boundaries, and defects using HAADF-STEM. Proc Natl Acad Sci 2010;107(14):6152e7. [16] Kirkland J. Advanced computing in electron microscopy. SpringerLink: Springer e-Books, Springer London, Limited; 2010. [17] Allen LJ, Findlay SD, Oxley MP, Rossouw CJ. Lattice-resolution contrast from a focused coherent electron probe. Part i. Ultramicroscopy 2003;96(1):47e63. [18] LeBeau JM, Findlay SD, Allen LJ, Stemmer S. Quantitative atomic resolution scanning transmission electron microscopy. Phys Rev Lett 2008;100:206101. [19] De Graef M. Introduction to conventional transmission electron microscopy. Cambridge University Press; 2003. [20] Gonzalez R, Woods R, Eddins S. Digital image processing using MATLAB. 2nd ed. NJ: Pearson Prentice Hall; 2009 [Chapter 6]. [21] Bolívar Guarín ME, Moreira A d M, Speziali NL. Structural study of Yb1xTbxMnO3 manganites: site occupancy and coexistence of hexagonal and orthorhombic phases. J Appl Crystallogr 2013;46(3):644e8. [22] Nascimento FC, Santos A O d, Campos A d, Gama SA, Cardoso LP. Structural and magnetic study of the MnAs magnetocaloric compound. Mater Res 2006;9:111e4. [23] Lin C, Liu J, Li X, Li Y, Chu S, Xiong L, et al. Phase transformation in hexagonal ErMnO3 under high pressure. J Appl Phys 2012;112(11):113512. [24] Kitano Y, Komura Y, Kajiwara H, Watanabe E. Two-dimensional lattice image of the Mg-base Friauf-Laves phase and a new type of defect. Acta Crystallogr Sect A 1980;36(1):16e21. [25] Komura Y, Takeda S, Takata M. Defect structures in the Mg-base Laves phase Mg (Cu0.08Zn0.92)2. Trans Jpn Inst Met 1983;24(6):413. [26] Takeda S, Kitano Y, Komura Y. Observations of the defects in the Laves phase Mg(Cu0.10Zn0.90)2 by high resolution electron microscopy. J Electron Microsc 1983;32(2):105e14. [27] Stein F, Palm M, Sauthoff G. Structure and stability of Laves phases. Part i. Critical assessment of factors controlling Laves phase stability. Intermetallics 2004;12(7e9):713e20. [28] Stein F, Palm M, Sauthoff G. Structure and stability of Laves phases. Part ii. Structure type variations in binary and ternary systems. Intermetallics 2005;13(10):1056e74.
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Intermetallics journal homepage: www.elsevier.com/locate/intermet
Phase coexistence in Ti6Sn5 intermetallics A.A. Oni, D. Hook, J.P. Maria, J.M. LeBeau* Department of Materials Science & Engineering, North Carolina State University, Raleigh, NC 27695-7907, USA
a r t i c l e i n f o
a b s t r a c t
Article history: Received 23 December 2013 Received in revised form 7 March 2014 Accepted 10 March 2014 Available online
Here we report the structural characterization of a complex TieSn intermetallic compound, Ti6Sn5. From X-ray diffraction, the resulting compound was observed to exist in both orthorhombic and hexagonal phases. Analysis by electron microscopy revealed that “planar-like” defects form throughout the material. Atomic resolution aberration-corrected scanning transmission electron microscopy reveals that these “planar-like” defects represent the coexistence of the orthorhombic and hexagonal phases within single grains. The resulting interwoven phases range in thickness from a fraction to multiple unit cells and exhibit coherent phase boundaries with the matrix grain. Ó 2014 Elsevier Ltd. All rights reserved.
Keywords: A. Intermetallics, miscellaneous D. Phase interfaces D. Defects: planar faults F. Electron microscopy, transmission F. Diffraction
1. Introduction Titanium alloys and intermetallics are utilized in many applications spanning aircraft production, medical instruments [1] and mechanical dampening [2]. They exhibit outstanding oxidation and corrosion resistance, low density, and mechanical properties [3,4]. Further, the electrical and magnetic properties of the intermetallics offer the potential for additional functionality. Experimental studies of hexagonal Ti6Sn5, for example, revealed a ground state in close proximity to a non-magnetic-magnetic phase boundary [5]. Moreover, doping Ti6Sn5 with rare earth elements introduces ferromagnetic instability and ordering [6]. Further, little is known about defects in Ti6Sn5 compounds, which may ultimately have important consequences on properties. The complex Ti6Sn5 intermetallic line compound has been reported to exhibit hexagonal, h-Ti6Sn5 (space group P63/mmc) [3], and orthorhombic, o-Ti6Sn5 (space group Immm) [7], phases. The lattice parameters for h-Ti6Sn5 and o-Ti6Sn5 are a ¼ 0.9248 nm, c ¼ 0.5690 nm [3] and a ¼ 1.693 nm, b ¼ 0.9144 nm and c ¼ 0.5735 nm [7], respectively. The phase diagram of the TieSn system [8] shows that hexagonal Ti6Sn5 (h-Ti6Sn5) is energetically favorable above 751 C, and the o-Ti6Sn5 phase is stable below [5]. Furthermore, the enthalpy of formation, DHf reported from firstprinciple calculations [9,10] for h-Ti6Sn5 and o-Ti6Sn5 is 37.25 kJ/mol and 34.92 kJ/mol, respectively [11]. Given the
* Corresponding author. Tel.: þ1 919 515 5049. E-mail address: [email protected] (J.M. LeBeau). http://dx.doi.org/10.1016/j.intermet.2014.03.002 0966-9795/Ó 2014 Elsevier Ltd. All rights reserved.
negligible difference in DHf, either phase may form depending on processing conditions. For example, experiments have demonstrated that the h-Ti6Sn5 phase is typically accompanied by orthorhombic o-Ti6Sn5 [5]. Defect analysis of complex intermetallic structures is challenging due to their large lattice parameters (small diffraction vector spacing) and because they possess a large number of atoms per unit cell, ranging from many tens of atoms up to more than a thousand [12,13]. The introduction of aberration correction in scanning transmission electron microscopy (STEM) has opened the doorway to the characterization of defects in these structures with unprecedented clarity [12,14,15]. In particular, high-angle annular dark-field (HAADF) STEM has proven an invaluable tool for materials characterization as its image intensities scale with atomic number. For example, aberration-corrected STEM was used to investigate stacking faults and dislocation cores in the Laves phase Cr2Hf, as a result, improving our general understanding of deformation mechanism and phase transformation kinetics in complex structures [14]. In this Article, we present a detailed structural analysis of the Ti6Sn5 intermetallic compound. By conventional BF/DF TEM, we show the formation of “planar-like” defects formed throughout grains of Ti6Sn5, which are reminiscent of stacking faults. We employ aberration-corrected high-angle annular dark-field scanning transmission electron microscopy (HAADF STEM) to determine the atomic scale structure of these defects. Structural models and HAADF-STEM image simulations are used to confirm that these “planar-like” defects result from the coexistence of h-Ti6Sn5 and oTi6Sn5 phases within single grains.
A.A. Oni et al. / Intermetallics 51 (2014) 48e52
2. Materials and methods Bulk TieSn mixtures were prepared from w20 mm Ti and w5 mm Sn powders combined in a 6:5 atomic ratio. The mixture was then homogenized with ball milling for 15 min. The powder was then compressed in a 1 cm die at 12 kpsi. The pellet was then heated to 1275 C for 17 h in flowing Ar ðpO2 z105 atmÞ. X-ray diffraction experiments were performed with a PANalytical Empyrean diffractometer equipped with a variable-slit scanning detector and Cu Ka (l ¼ 1.5418 nm) incident radiation. Specimens for electron microscopy were prepared by conventional mechanical polishing using an Allied High Tech MultiprepÔ. A Fischione Model 1050 Ar ion mill with a LN2 cooled stage was used to thin samples to electron transparency. Conventional bright-/dark-field images were acquired with a JEOL 2000FX transmission electron microscope operated at 200 kV and equipped with a LaB6 emitter. A probe-corrected FEI Titan G2 60-300 KV S/ TEM equipped with a high-brightness Schottky field emission gun was used for HAADF-STEM imaging at 200 kV. The convergence and collection semi-angles for the Titan S/TEM were 21 mrad and 76e 400 mrad respectively. Atomic resolution HAADF-STEM image simulations were performed using the multislice approach [16]. The images were simulated using the following input parameters: sample thickness ¼ 20 nm, C3 ¼ 6 mm, C5 ¼ 2.5 mm, and the same convergence and collection geometries that were used during experimental micrograph acquisition. Thermal diffuse scattering was introduced in the simulations using the frozen lattice model [17] and an estimated root-mean-squared (r.m.s.) displacement of 0.09 A for both Ti and Sn. Simulated images were convolved with a Gaussian function, 0.1 nm full-width half-max, to approximately account for the finite effective source size [18]. Simulated images of phase coexistence regions were generated by joining separate simulations of h-Ti6Sn5 and o-Ti6Sn5 for the same crystal thickness. 3. Results and discussion X-ray diffraction analysis of the as-processed bulk pellet, presented in Fig. 1, showed the formation of the Ti6Sn5 intermetallic with trace amounts of both Ti5Sn3 and Ti2Sn5. The peaks labeled ‘Other’ correspond to Ti5Sn3 and Ti2Sn5. The presence of the hTi6Sn5 and the o-Ti6Sn5 phases was confirmed by atomic resolution HAADF-STEM imaging as shown in Fig. 2. The h-Ti6Sn5 phase was the most frequently observed throughout the bulk of the specimen. The crystal structure of h-Ti6Sn5 consists of ABCB0 ABCB0 stacking sequence when viewed along [100] projection as seen in Fig. 2(a).
Fig. 1. Powder X-ray diffraction pattern of the TiSn alloy. h-Ti6Sn5 and o-Ti6Sn5 peaks are indicated, while traces of Ti2Sn5 and Ti5Sn3 phases detected are presented as Other phases.
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The A planes consist of atom columns containing both Sn atoms and Ti atoms. The B ‘planes’ exhibit alternating atom column doublets (TieSn) and triplets (TieSneSn). The C plane contains only Ti atom atoms. The B0 and B planes are related through a two-fold rotation about the C plane. The stacking sequence of the o-Ti6Sn5 phase along the [010] projection is nearly identical to h-Ti6Sn5 except that mirror symmetry is introduced across the A planes, such that the sequence becomes ABCB0 AB0 CB. Comparison of HAADF-STEM images from experiment and simulations along the hexagonal h100ih (Fig. 2(b)) shows the closely-spaced SneTi (doublet) columns have the brightest intensity and the pure Ti columns have the dimmest intensity. Furthermore, a similar intensity variation in atomic columns is noted in o-Ti6Sn5, in Fig. 2(c), however, there are important differences between the HAADF-STEM images of the two phases. A chevron-like pattern is observed for the hexagonal phase, as noted with white lines in Fig. 2(b), whereas mirror symmetry is present across the C planes of o-Ti6Sn5 (ABCB0 AB0 CB sequence). These image intensity details are essential for direct identification of the crystal structure and defect structures in this system. At the microstructural level, conventional transmission electron microscopy revealed planar defects throughout the Ti6Sn5 grains, as indicated in Fig. 3(a). Individual grains typically possessed a high density of these planar defects, with roughly 0.10 per nm. Furthermore, selected area diffraction patterns along [010] showed streaking of the diffraction spots parallel to g100, see Fig. 3(b). The observed streaks are indicative of randomly displaced (100) planar faults, and are typically observed with high stacking fault or twin boundary density structures [19]. Further, the planar-type defects were investigated with aberration-corrected STEM, where they appear as a disturbance in the stacking sequence and a corresponding decrease in the image intensity, as indicated by the white arrows in Fig. 3(c). Close examination of the atomic structure in Fig. 3(c) suggests the possibility that these defects are the formation of 180 domain boundaries, where the chevrons of the [100]h projections mirror across the defect plane relative to one another. Rather than abrupt boundaries, however, the defects exhibited nm-scale spatial extent perpendicular to the (100)h fault plane ranging from approximately 1e10 nm. Moreover, in regions where the defects appeared to extend beyond a single layer, atomic resolution HAADF imaging revealed the formation of multiple unit cells of o-Ti6Sn5 within an h-Ti6Sn5 matrix. An example of two unit cells of h-Ti6Sn5 embedded within an o-Ti6Sn5 is shown in the HAADF-STEM image in Fig. 3(d). These coexistence defects are consistent with the interpretation of o-Ti6Sn5 within an h-Ti6Sn5 matrix. From inspection of Fig. 3(c) and (d), image simulations qualitatively reproduce the features across the boundary. An averaged intensity line profile from experiment is presented in Fig. 4 and compared to the simulation. To increase the signal-to-noise ratio, the boxed region indicated in Fig. 3(c) was used as a template to find identical phase coexistence regions within the experimental image using a normalized crosscorrelation approach [20]. These regions were subsequently averaged and integrated along [001]h to generate the line profile. As shown, the image contrast is also in good qualitative agreement when employing the phase coexistence model. To complement the results from [100]h, the planar defects were also observed along [011]h, as seen in Fig. 5. The defects again extended beyond a single plane and the structure at the boundary was inconsistent with anti-phase or domain boundaries, indicating the need for an alternative structural model to explain the defect features. Rather, the introduction of a sub-unit layer or layers of oTi6Sn5 provides the necessary symmetry needed to account for the defect mirror plane and spatial extent. The phase coexistence model is shown schematically in Fig. 5. As indicated in the stacking
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Fig. 2. (a) HAADF-STEM image of h-Ti6Sn5 in the h100i projection. (b) HAADF-STEM image of o-Ti6Sn5 in the h010i projection. (c) Schematic showing stacking sequence of h-Ti6Sn5 and o-Ti6Sn5 structure. Atomic model and HAADF simulations are superimposed on the HAADF-STEM images.
sequence schematic, the stacking sequence for h-Ti6Sn5 in this projection is DED0 E0 D, where D0 and E0 are related to D and E through inversion symmetry across the sub-unit boundary. For comparison, the stacking sequence for o-Ti6Sn5 is simply DEDE. Along the [011] zone, the atom columns project a distinctly different atom positions along [011]h and [011]o. The formation of these defects in the Ti6Sn5 correlates with the minimal difference in DHf for h-Ti6Sn5 and o-Ti6Sn5, and are likely to result from processing conditions. As a result, these defects may not occur in all cases and will depend upon the kinetics in addition to the thermodynamic stability. Moreover, these results indicate the possibility that, rather than the formation of stacking faults or other planar defects, the coexistence of h-Ti6Sn5 and o-Ti6Sn5 acts
to relieve stress during mechanical deformation. Furthermore, given the close proximity of the hexagonal phase to non-magneticmagnetic phase boundary, these results may have important consequences on controlling the magnetic properties of Ti6Sn5 intermetallic compounds. Phase coexistence of hexagonal and orthorhombic phases has also been observed by varying composition [21], temperature [22], and pressure [23]. For example, structural defects arising from stacking variants have been observed in the Mg-based pseudo-binary Laves phases [24e26]. As we have observed for Ti6Sn5, Laves phases are capable of coexisting in the same system where they exhibit abrupt changes in the stacking sequence [27,28] and similar to Ti6Sn5, the Laves phases exhibit fundamental prototypes with a
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Fig. 3. (a) Two-beam dark-field TEM imaging of a Ti6Sn5 grain using g100 reflection. Planar defects are indicated with arrows. (b) [010] zone axis selected area diffraction pattern. The streaking of the diffraction spots are indicative of high stacking fault faults or twin boundary. (c) HAADF-STEM image of h-Ti6Sn5 in the h100i projection with o-Ti6Sn5 phase embedded within the matrix. (d) h-Ti6Sn5 phase embedded within o-Ti6Sn5.
range of crystal structures, namely cubic e C14, and hexagonal e C15 and C36 [27]. 4. Conclusions Through a combination of X-ray diffraction, conventional TEM, and aberration-corrected STEM, we have demonstrated that the identification of defects in complex intermetallic structures
Fig. 4. Average intensity line profile extracted from an HAADF-STEM image of an hTi6Sn5eo-Ti6Sn5 phase coexistence region (black, solid line e see indicated area in Fig. 3(c)) compared with simulation (red, dashed line). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 5. HAADF images of phase coexistence of h-Ti6Sn5 and o-Ti6Sn5 in the h011i projection. The inset is a schematic for the stacking sequence of h-Ti6Sn5 and o-Ti6Sn5 in the h011i projection.
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remains a challenging endeavor. Specifically, we have shown that conventional TEM analysis can lead to misidentification of planar defects when phase coexistence occurs. In the case of Ti6Sn5 planar defects, the coexistence of hexagonal and orthorhombic phases is readily revealed by atomic resolution imaging. Moreover, our results demonstrate the power of aberration-corrected STEM to fully characterize structural defects in complex structures. Additional work is required to relate the observed phase coexistence with the mechanical and magnetic properties of Ti6Sn5. Acknowledgments The authors acknowledge Oak Ridge National Laboratory ShaRE program for the preliminary data acquisition for this project. A.A.O and J.M.L acknowledge support from the Air Force Office of Scientific Research (Grant no. FA9550-12-1-0456). J.P.M and D.H. acknowledge funding for this project through DuPont Microcircuit Materials under contract 2008-2446. The authors acknowledge the use of the Analytical Instrumentation Facility (AIF) at North Carolina State University, which is supported by the State of North Carolina and the National Science Foundation. References [1] Wang X, Li W, Fang G, Wu C, Lin J. First-principles calculations on the electronic structure and cohesive properties of titanium stannides. Intermetallics 2009;17(9):768e73. [2] Wong CR, Fleischer RL. Low frequency damping and ultrasonic attenuation in Ti3Sn-based alloyssed alloys. J Mater Res 1994;9(6):1441e8. [3] Villars P, Calvert LD. Pearson’s handbook crystallographic data for intermetallic phases, vol. 2. Metals Park, OH: American Society for Metals; 1985. [4] Kim Y-W, Dimiduk D. Progress in the understanding of gamma titanium aluminides. JOM 1991;43(8):40e7. [5] Drymiotis F, Lashley JC, Fisk Z, Peterson E, Nakatsuji S. Physical properties of the b-Ti6Sn5 system. Philos Mag 2003;83(27):3169e78. [6] Jeong T. Electronic structure and magnetic properties of b-Ti6Sn5. J Magnet Magnet Mater 2007;309(1):71e4. [7] van Vucht JHN, Bruning HACM, Donkersloot HC, de Mesquita AHG. Philips Res Rep 19(5);407e21. [8] Okamoto H. SneTi (TineTitanium). J Phase Equilibria Diff 2010;31(2):202e3.
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