Article pubs.acs.org/cm
Lithium Atom and A‑Site Vacancy Distributions in Lanthanum Lithium Titanate Xiang Gao,† Craig A. J. Fisher,† Teiichi Kimura,† Yumi H. Ikuhara,† Hiroki Moriwake,† Akihide Kuwabara,† Hideki Oki,‡,§ Takeshi Tojigamori,‡ Rong Huang,¶ and Yuichi Ikuhara*,†,# †
Nanostructures Research Laboratory, Japan Fine Ceramics Center, Nagoya 456-8587, Japan Battery Materials Division, Toyota Motor Corporation, Shizuoka 410-1193, Japan § Graduate School of Frontier Science, The University of Tokyo, Tokyo 113-0033, Japan ¶ Key Laboratory of Polar Materials and Devices, Ministry of Education, East China Normal University, Shanghai 200062, China # Institute of Engineering Innovation, The University of Tokyo, Tokyo 113-8586, Japan ‡
S Supporting Information *
ABSTRACT: Lanthanum lithium titanate (LLTO) is one of the most promising electrolyte materials for all-solid-state lithium-ion batteries. Despite numerous studies, the detailed crystal structure is still open to conjecture because of the difficulty of identifying precisely the positions of Li atoms and the distribution of intrinsic cation vacancies. Here we use subangstrom resolution scanning transmission electron microscopy (STEM) imaging methods and spatially resolved electron energy loss spectroscopy (EELS) analysis to examine the local atomic structure of LLTO. Direct annular bright-field (ABF) observations show Li locations on O4 window positions in Li-poor phase La0.62Li0.16TiO3 and near to A-site positions in Li-rich phase La0.56Li0.33TiO3. Local clustering of A-site vacancies results in aggregation of Li atoms, enhanced octahedral tilting and distortion, formation of O vacancies, and partial Ti4+ reduction. The results suggest local LLTO structures depend on a balance between the distribution of A-site vacancies and the need to maintain interlayer charge neutrality. The associated local clustering of A-site vacancies and aggregation of Li atoms is expected to affect the Li-ion migration pathways, which change from two-dimensional in Li-poor LLTO to three-dimensional in Li-rich LLTO. This study demonstrates how a combination of advanced STEM and EELS analysis can provide critical insights into the atomic structure and crystal chemistry of solid ionic conductors. KEYWORDS: solid-state electrolyte, ionic conductor, structure−property relationships, Li-ion battery, scanning transmission electron microscopy
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INTRODUCTION The continuing drive for high performance lithium batteries to power fully electric and hybrid vehicles imposes stricter demands on the electrolyte materials they utilize.1−3 Many of these demands can be met by all-solid-state lithium-ion batteries containing nonflammable solid electrolytes, which have important advantages over conventional commercial (liquid-electrolyte- or polymer-electrolyte-based) batteries such as greater thermal stability, lack of electrolyte leakage, increased cycle life and energy density, and robustness to vibration and physical impacts.4−7 For these advantages to be exploited, a solid electrolyte is needed that has both high lithium-ion conductivity and good chemical stability. The superionic conductor La2/3‑xLi3xTiO3 (LLTO) has been reported to have a high bulk ionic conductivity of up to 1.0 × 10−3 S cm−1 (for x ≈ 0.11) at room temperature,and is one of the most promising candidate materials for all-solid-state lithium-ion batteries.7,8 Much effort has been expended in attempts to identify which structural features enhance Li-ion mobility in this material. Powder diffraction studies, either using X-rays (XRD) or © XXXX American Chemical Society
neutrons (ND), have shown that LLTO has an ABO3 perovskite-type structure but one in which the A sites contain a large number of defects. Indeed, the end member La2/3TiO3 is intrinsically A-cation deficient, with one-third of A sites vacant.9 Rather than being randomly distributed, at normal temperatures these La vacancies are partitioned into alternating La-rich and La-poor layers along one axis to form a partially ordered superlattice structure (Figure 1).10−19 Several studies have examined the variation of the crystal structure of LLTO with Li content, with sometimes conflicting results. Li-poor (0.03 ≤ x < 0.1) compositions are generally reported to exhibit orthorhombic symmetry, with high La-site occupancies (≥90%) in the La-rich layer together with antiphase tilting of Ti−O6 octahedra;10−14 in Li-rich (0.1 ≤ x < 0.167) compositions, the symmetry is tetragonal, and the occupancies of the two types of La layers become less dissimilar as the Li content increases.13−19 Received: December 27, 2012 Revised: April 8, 2013
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composition variations are small. Such an advantage is best realized with a spherical-aberration corrected scanning transmission electron microscope (STEM), which uses a subangstrom electron probe. Recently, atomic resolution annular bright-field (ABF) STEM imaging methods have been shown not only to enable simultaneous imaging of columns of both light and heavy atoms but also to enable robust interpretation over a wide range of specimen thicknesses.31,32 This technique has been successfully used to image H (Z = 1) and Li (Z = 3) atoms, the lightest and third lightest elements in the periodic table, respectively, directly in real space and real time.33−36 Simultaneous HAADF and ABF imaging in an aberration-corrected STEM is thus a powerful means of observing subtle structural and chemical variations in a crystal at the atomic scale, even for relatively complex crystal structures such as in LLTO.37 In this paper, we report the first STEM study of the atomic structure and crystal chemistry of LLTO perovskites using combined direct HAADF/ABF imaging methods and spatially resolved electron energy loss spectroscopy (EELS) analysis in a spherical-aberration corrected STEM. Two LLTO compositions are examined: a Li-poor composition, La0.62Li0.16TiO3, and a Li-rich composition, La0.56Li0.33TiO3.
Figure 1. Perspective views of the LLTO perovskite network along [100]p and [110]p zone axes (where ‘p’ refers to the cubic pseudoperovskite structure).20 Li atoms are not shown on account of uncertainties regarding their positions reported in the literature.11,12,15−17,21−23 O4 square window sites are indicated by dotted circles.
Other symmetries can be obtained through quenching.24 Notably, XRD measurements of La0.5Li0.5TiO3 (x = 0.167) quenched from high temperature indicated a cubic structure with La cations randomly distributed over equivalent A sites,13,25 although ND and Rietveld analysis have since revealed the structure is actually rhombohedral.26 Despite these numerous investigations, the precise crystal structure of LLTO, particularly the locations of Li ions, is still a matter of controversy. It became clear early on that the small Li cations do not sit on the A sites themselves; off-center positions next to vacant La sites and O4 square window positions midway between La sites in La-poor layers have both been proposed as the location of Li within the Li-poor perovskite La0.62Li0.16TiO3.11,12 Li-rich perovskites La0.56Li0.33TiO3 and La0.567Li0.3TiO3, corresponding to x ≈ 0.11, were reported to have the same tetragonal structure, but with different space group (S.G.) symmetries of P4/mmm and P4/nbm,15,17 respectively. In the latter case, Li was reported to be located on O4 square window sites. In contrast, La0.55Li0.35TiO3 has been reported to be orthorhombic (S.G. Cmmm), with Li occupying both the O4 square window (Wyckoff 2c) positions in La-rich layers and off-center positions of vacant La sites in La-poor layers.16 Theoretical calculations based on quantummechanical methods, classical molecular dynamics simulations, and bond valence summation have also found different optimal Li positions for various LLTO compounds.21−23,27 High-resolution electron microscopy (HREM) has previously revealed that these materials have complex microstructures composed of microdomains with different periodicities and crystal orientations.15,24,28 It has been reported that at low Li contents the maximum domain size is around 1000 Å2, while for high Li contents it is about 10 times smaller.15 The reasons for these microstructural changes are still poorly understood. Given the complexity of this highly defective system, it is not surprising then that conflicting Li diffusion mechanisms have been proposed to explain the rapid Li-ion conductivity. One way to help resolve this issue is to observe the atom positions and crystal structure of this superionic conductor material directly. High-angle annular dark-field (HAADF) imaging provides directly and robustly interpretable images of the crystal structure, with contrast roughly proportional to the square of the atomic number (Z) of the atoms being imaged,29,30 making it possible to determine the distribution of all but the lightest elements in a material at the atomic level even when local
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EXPERIMENTAL SECTION
Specimen Preparation. Ceramic Li-poor (La0.62Li0.16TiO3) and Li-rich (La0.56Li0.33TiO3) samples were prepared by a solid-state reaction method from mixtures of Li2CO3 (Wako Pure Chemical Industries, > 99%), TiO2 (Wako Pure Chemical Industries, > 99.9%), and freshly dehydrated La2O3 (Wako Pure Chemical Industries, > 99.99%). An excess of 20 mol % Li2CO3 was added to the stoichiometric mixtures to compensate for Li loss during processing. The mixed powders were calcined at 800 °C for 8 h and pressed into pellets before sintering at 1250 °C for 12 h, after which they were furnace-cooled to room temperature. Pellets were embedded in the remaining calcined powder in a covered Al2O3 crucible. Crystalline phase identification was carried out by XRD with Cu-Kα radiation using a Rint2000 diffractometer (Rigaku, Tokyo, Japan). The compositions of sintered samples were analyzed by inductive coupled plasma (ICP) spectroscopy. Specimens for TEM and STEM observations were prepared by cutting, grinding, dimpling, and ionmilling (Gatan 691), operated at 3.0 to 0.5 kV and cooled with liquid nitrogen, followed by surface cleaning using an ion cleaner (JIC-410, JEOL, Japan) to remove completely any residual amorphous film. Microscopy. A 300 kV JEM-3000F (JEOL) microscope was used for HREM observations. HAADF/ABF-STEM observations were performed using a 200 kV JEM-2100F (JEOL) microscope equipped with a spherical-aberration corrector (CEOS GmbH) and a Gatan Image Filter (GIF) for EELS analysis, which enabled structures to be probed with subangstrom resolution. Z-contrast HAADF imaging was performed in thin specimen regions with a probe convergence angle of 25 mrad and an annular dark-field detector inner angle greater than 52 mrad. ABF images were taken simultaneously using a 9−19 mrad detector, and the EELS spectra were recorded with a collection semiangle of 35 mrad. All HAADF and ABF images were recorded using the same dwell time of 31 μs. ABF-STEM image simulation was carried out with the multislice method, using experimentally measured values for lens and detector aperture settings. Note that the specimen thicknesses assumed may not correspond to those in the experiment.
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RESULTS AND DISCUSSION XRD analysis confirmed that crystalline LLTO was obtained as a pure phase for both the Li-rich and Li-poor samples. In agreement with previous reports, the Li-rich sample had a tetragonal crystal structure, while the Li-poor sample was B
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orthorhombic, as indicated by splitting of the (200) peak13 (see the Supporting Information). ICP analysis also revealed that the elemental ratios in the two samples were close to their nominal values, with only a small deviation in the Li content as a result of Li loss during processing. The measured compositions were found to be Li0.12La0.61TiO3‑d and Li0.30La0.54TiO3‑d for the Li-poor and Lirich samples, respectively (compared to nominal compositions of Li0.16La0.62TiO3 and Li0.33La0.56TiO3, respectively). Figure 2a shows elongated domains with straight 90°oriented domain boundaries aligned parallel to [110]P (where p indicates the primitive perovskite unit cell) in the Li-poor sample La0.62Li0.16TiO3. In contrast, the Li-rich sample (La0.56Li0.33TiO3) in Figure 2b exhibits a mosaic-like structure comprising smaller domains. This dependence of domain
structure on Li content is consistent with previous HREM observations.15,24,28 An important common feature is noticeable in both cases: within an individual domain, atom layers alternate between rows of uniformly bright dots and rows of more or less dark dots (examples are arrowed in white and black, respectively, in Figure 2) along the c axis. This stark variation corresponds to the ordering of La in alternate (001) layers. The superstructure reflection of 1/2(001) in the corresponding selected area electron diffraction (SAED) patterns (insets of Figure 2a,b) further confirms the ordering along c. It should be noted here that the diffraction patterns of the Li-rich sample always include several domains, which are too small (
