Epitaxial Growth of LiMn2O4 Thin Films by Chemical Solution

Aug 5, 2014 - Toyota Motor Corporation, Susono, Shizuoka 410-1193, Japan. •S Supporting ... Transmission electron microscopy (TEM) and scanning...
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Epitaxial Growth of LiMn2O4 Thin Films by Chemical Solution Deposition for Multilayer Lithium-Ion Batteries Yumi H. Ikuhara,*,† Xiang Gao,† Rong Huang,†,‡ Craig A. J. Fisher,† Akihide Kuwabara,† Hiroki Moriwake,† and Keiichi Kohama§ †

Nanostructures Research Laboratory, Japan Fine Ceramics Center, 2-4-1 Mutsuno, Nagoya 456-8587, Japan Key Laboratory of Polarized Materials and Devices, East China Normal University, Shanghai, 200241, China § Toyota Motor Corporation, Susono, Shizuoka 410-1193, Japan ‡

S Supporting Information *

ABSTRACT: Cathodic LiMn2O4 films on various single-crystal substrates for use in thin-film Li-ion batteries prepared using a chemical solution deposition method are reported. Transmission electron microscopy is utilized to characterize the microstructures of the films. The results show that the film/substrate lattice misfit can affect significantly the quality of epitaxially grown grains of LiMn2O4. Using state-of-the-art high-angle annular dark-field imaging, the degree of coherency and lattice distortion at interfaces between LiMn2O4 and Au-coated and uncoated Al2O3 (0001) single-crystal substrates are examined at the atomic scale. When the lattice misfit is sufficiently small, fully coherent LiMn2O4/Au heterointerfaces form, although lattice strain to a distance of up to around 10 nm from the interface changes the symmetry of spinel LiMn2O4 from cubic to tetragonal. Such an interface in the LiMn2O4/Au/Al2O3 system facilitates high-quality epitaxial film growth to thicknesses of a couple hundred nanometers.

deposition,15,16 and chemical processing via a solution route17,18 to produce small-, medium-, and large-scale batteries as required. Chemical solution deposition (CSD) using metalorganic precursors is particularly attractive for fabricating thin films because it not only provides close control of the material’s purity and compositional homogeneity with desired stoichiometry,19,20 but the growth rate is also comparatively high, and it does not need to be performed under vacuum. All-solid-state batteries present a different set of challenges to conventional liquid-electrolyte batteries, particularly in regard to increased interfacial resistances and detrimental secondary phase formation between the various component layers.21 A proper understanding of the microstructure of the film and interfaces both within the film and between the film and substrate is needed to optimize battery performance using nanoengineering techniques, because electrochemical properties can vary significantly with the crystal orientation, the degree of lattice strain, and the nature of the terminating layers.22−25 For example, it has been shown that the LiMn2O4 (111) surface is more stable than the LiMn2O4 (110) surface upon formation of the solid electrolyte interface layer when in contact with a liquid electrolyte during electrochemical measurements.12 In

INTRODUCTION Rechargeable lithium-ion secondary batteries are being developed for use in high-power applications such as fully electric and hybrid electric vehicles because of their high energy and power densities compared to other battery technologies.1−3 Lithium manganese oxide (LiMn2O4, or LMO) is an attractive cathode material for such batteries because of its low cost (given the relative abundance of Mn), low toxicity, greater safety, and environmental compatibility compared to the conventional layered cathode material LiCoO2.4−6 Its crystal structure is of the spinel type (space group Fd3m), with lithium ions occupying one-eighth of the tetrahedral interstices within a cubic-close-packed array of oxide ions and manganese ions on half the octahedral interstices. In contrast to other well-known cathode materials such as LiCoO2 and LiFePO4, the threedimensional network between tetrahedral sites enables lithium ions to be reversibly intercalated into the structure in all three crystallographic directions.4,7 All-solid-state batteries comprising multilayer thin films are being actively developed for the above-mentioned applications because they remove the need to use a toxic, flammable, or caustic liquid electrolyte, improving battery safety, while also increasing the battery capacity and energy density because less packaging material is required.8 Suitable LiMn2O4 thin films for this can be synthesized by a variety of techniques, including sputtering,9,10 pulsed laser deposition,11−14 electrostatic spray © 2014 American Chemical Society

Received: May 2, 2014 Revised: August 2, 2014 Published: August 5, 2014 19540

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the case of all-solid-state batteries, one strategy for improved design and performance of multilayered thin films is to form an epitaxial cathode film, so that preferred Li intercalation/ deintercalation pathways can be controlled and maintained during the lifetime of the battery.26,27 Transmission electron microscopy (TEM) and scanning transmission electron microscopy (STEM) are powerful tools for examining crystalline bulk and interfacial structures. In particular, microscopes equipped with state-of-the-art aberration correctors now allow crystal structures to be imaged directly with subangstrom resolution. High-angle annular dark field (HAADF) detectors enable the positions of relatively heavy elements to be viewed in real time, allowing the orientation relationship between film and substrate and the details of atomic alignments and configurations to be determined with unprecedented clarity. The increasing importance of these improved techniques for the direct visualization of crystal structures, as well as the ability to identify point defects, local clusters, and extended interfaces at the atomic level, is illustrated in the case of lithium battery materials by recent studies on nanoparticles of LiMn2O4,28,29 phase development in thin films of LiCoO2,30 cation-vacancy distributions and domain boundary structures in La 2/3−xLix TiO3 ,31,32 and changes during cycling in LiNi0.8Co0.15Al0.05O2.33 The present study aims to investigate the effects of lattice mismatch on LiMn2O4 film growth behavior during CSD by comparing results for various substrates, namely Al2O3 (0001), MgO (110), SrTiO3 (110), and gold-coated single-crystal Al2O3 (0001) substrates. Since Al2O3, MgO, and SrTiO3 are electrically insulating, they are not themselves being considered as current collectors for the lithium-ion batteries. If such materials were to be used, an electrically conducting layer such as Au would need to be deposited before formation of the LiMn2O4 thin film, as was done in the case of the LMO/Au/ Al2O3 sample. After outlining the synthesis and characterization methods in the next section, a detailed examination of the thin films using TEM and STEM is presented in the third section in an effort to understand the relationship between substrate surface orientations, lattice misfit strain, and the quality of thin films that can be produced.

Figure 1. Flowchart showing steps in the preparation of LiMn2O4 thin films on single-crystal substrates by the chemical solution deposition (CSD) method.

(110), and uncoated and Au-coated Al2O3 (0001) single crystals. The fourth substrate was prepared by depositing Au films on Al2O3 (0001) single crystals by rf-diode sputtering. The target was a 99.99% Au plate that covered the cathode of the rf sputterer. Before sputtering, Al2O3 (0001) substrates were chemically cleaned and set on a table which covered the anode. The chamber was evacuated to a pressure of 3 × 10−4 Pa and backfilled with 99.999% pure argon gas. The argon gas pressure was kept constant at 1.5 Pa throughout the deposition, the substrate temperature was held at 650 °C, and the sputtering power was set at 100 W. LiMn2O4 precursor films were fabricated from the browncolored [Li−Mn−O] metalorganic precursor solution by spin coating onto the various substrates at a spinning speed of 2000 rpm for 20 s under flowing nitrogen gas. The LiMn2O4 precursor films were then transferred to a single-zone tube furnace and heated at 200 °C for 30 min to remove the organic ligands, before being heated to either 500, 600, or 750 °C at a rate of 5 °C/min, followed by annealing for 1 h in oxygen. To obtain thicknesses greater than 100 nm, this procedure was repeated several times. X-ray diffraction (XRD) measurements of the heat-treated films were performed with Cu Kα radiation using a Rint2000 (Rigaku, Tokyo, Japan) diffractometer equipped with a monochromator, operating at 40 kV and 50 mA. The morphologies, orientation relationships, and detailed interface structures of the thin films after heat treatment were characterized by TEM and STEM. Samples for cross-sectional observations were prepared as follows. Each sample was cut using a diamond saw parallel to the desired cross-sectional plane. Next, for each system, two cut pieces were glued together with the film surfaces facing each other. These samples were cut into lengths of about 1.5 mm, polished to a thickness of ≈0.1 mm, and attached to molybdenum rings for reinforcement. These samples were subsequently dimpled and finally thinned by ion-beam sputtering at a voltage of 5 kV, using a cold stage (liquid nitrogen) to minimize damage to the specimens. TEM observations were performed using a JEOL 2010 electron microscope (JEOL) operating at 200 kV with a pointto-point resolution of about 0.194 nm. Z-contrast HAADF observations were made using a 200 kV JEM-2100F (JEOL) scanning transmission electron microscope equipped with a spherical-aberration corrector (CEOS GmbH), providing a minimum probe of about 0.08 nm in diameter. During

EXPERIMENTAL METHODS Figure 1 shows the processing route to fabricate LiMn2O4/ substrate bilayer films by CSD. Metalorganic precursors for the fabrication of LiMn2O4 were LiOCH(CH3)2 and Mn(OC3H7)2 with 2-ethoxyethanol (C2H5OC2H4OH or EGMEE) as the solvent. Chemical modification of the precursors was carried out under dry nitrogen atmosphere. The precursors were dissolved in EGMEE and refluxed at 135 °C for 3 h. The alkoxy ligands in LiOCH(CH3)2 and Mn(OC3H7)2 were modified to give LiOC2H4OC2H5 and Mn(OC2H4OC2H5)4 by complete reaction. After cooling to room temperature, these two solutions were condensed in a rotary evaporator and added to fresh EGMEE solution to remove the subreacted products of (CH3)2CHOH. Lastly, a homogeneous metalorganic [Li−Mn− O] precursor solution was prepared by mixing the two prepared solutions of LiOC 2 H 4 OC 2 H 5 in EGMEE and Mn(OC2H4OC2H5)4 in EGMEE, followed by refluxing at 135 °C for 1 h. As mentioned in the introduction, four different substrate materials are compared in this study: SrTiO3 (110), MgO 19541

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HAADF-STEM imaging, a probe convergence angle of 25 mrad and an annular dark-field detector with an inner angle greater than 50 mrad were used. No filtering was performed on the HAADF-STEM images. Simulated HAADF-STEM images of the interface structure of LiMn2O4/Al2O3 were produced using CrystalKit and MacTempas software to aid interpretation of the observed structures. Cyclic voltammetry measurements were performed on the epitaxial cathode film deposited on Au-coated sapphire to assess the electrochemical properties of the film. A cell was made using lithium metal as the anode and a mixture of PC and 1 M LiClO4 as the electrolytic solution. A scan rate of 0.1 mV s−1 between 3.0 and 4.3 V was used.

RESULTS AND DISCUSSION To determine the optimum conditions for synthesis of the films, the temperature dependence of the growth of crystalline LiMn2O4 on a single-crystal Al2O3 substrate was first examined. In the prepared [Li−Mn−O] metalorganic precursor solution, Mn(OC2H4OC2H5)4 and LiOC2H4OC2H5 are expected to form large complexes, with each cation in a suitable coordination environment to react at elevated temperature in the desired molar ratio. In the case of LiMn2O4, nucleation is known to occur at 250 °C after evaporation of organic solvent without any second phase formation.18 This was found to be the case for thin-film formation of [Li−O−Mn] precursor films on an Al2O3 substrate also, as randomly oriented LiMn2O4 particles formed upon heat treatment at 250 °C. Based on an earlier study of the relationship between annealing temperature and crystallization of LMO particles,18 we estimate that the internal lattice strain in the nanoparticles was about 0.6%. Figure 2 shows cross-sectional bright-field TEM (BF-TEM) images and corresponding selected area electron diffraction (SAED) patterns of LMO thin films on Al2O3 (0001) substrates annealed at 500, 600, and 750 °C. The incident electron beam in the SAED patterns is parallel to the [112̅0]Al2O3 direction. LiMn2O4 (111) and (311) rings can be seen in the SAED patterns in Figure 2b and d, indicating that polycrystalline LMO films formed during heat treatment at 500 and 600 °C. Based on analysis of the cross-sectional BF-TEM images, randomly oriented small nuclei of LiMn2O4 about 5 nm in size were formed at 500 °C (Figure 2a). When the annealing temperature was increased to 600 °C, grain growth accelerated and the grain size of LMO particles in the film increased to about 20 nm. (Figure 2c). However, epitaxial LiMn2O4 grains did not develop. This may be because nanocrystalline LiMn2O4 particles heat treated at 500 and 600 °C still contain around 0.3% internal lattice strain,18 which, despite being relatively small, appears to inhibit the ability of the nuclei to form coherent bonds with the sapphire (0001) substrate. In contrast, Figures 2e and f show that increasing the annealing temperature to 750 °C produces a single-crystal LMO film, with orientation relationships [112̅]LiMn2O4//[112̅0]Al2O3 and (111)LiMn2O4// (0001)Al2O3. At this temperature, the nanoparticles are estimated to contain only around 0.1% internal lattice strain,18 with lattice parameters very close to those of the bulk material. These results show that annealing at 750 °C is sufficient to form epitaxial LiMn2O4 films on single-crystal Al2O3 substrates. Subsequent spin coating of the LMO precursor solution onto the other substrates, followed by heat treatment at 200 °C to evaporate the organic component, was thus followed by annealing at 750 °C for 1 h. The influence of each substrate

Figure 2. Cross-sectional bright-field TEM (BF-TEM) images and corresponding selected area electron diffraction (SAED) patterns showing the growth of LiMn2O4 thin films on Al2O3 (0001) substrates fabricated at (a and b) 500 °C; (c and d) 600 °C; and (e and f) 750 °C.

on epitaxial grain formation was then examined by XRD, TEM, and STEM techniques. The growth of epitaxial films is known to be controlled to a large extent by the lattice misfit parameter, f = (d1 − d2)/d2, where d1 and d2 correspond to the d-spacings between lattice planes in the expected epitaxial layer and substrate, respectively, parallel to the heterointerface. Table 1 summarizes calculated misfit parameters for cubic LiMn2O4 relative to the four substrates examined in this study, as well as the expected epifilm/substrate orientation relationship, crystal directions observed by cross-sectional TEM, and the inferred planes parallel to the interface. Lattice constants of 0.4078 nm,34 0.4211 nm,35 0.3905 nm,36 and 0.8244 nm28 for cubic Au, MgO, SrTiO3, and LiMn2O4, respectively, were used for the misfit calculations. The same calculations were performed using lattice constants for rhombohedral Al2O3 of a = b = 0.475 nm and c = 1.299 nm (JCPDS file #10-173). All the single crystals used in this study except Al2O3 have cubic symmetry. The calculated misfit parameters between (11̅0)LMO planes and parallel (11̅0)SrTiO3 and (11̅0)MgO planes, which are perpendicular to the interface, are 5.5% and −2.1%, respectively (Table 1). The smallest misfit parameter is 1.0% for an interface consisting of parallel (11̅0)LMO and (101̅)Au planes. Figure 3 shows cross-sectional BF-TEM images of LiMn2O4 thin films grown on each of the four substrates, together with their corresponding SAED patterns. In the case of LiMn2O4 on 19542

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Table 1. Theoretical Lattice Misfit Parameters of Likely Thin LiMn2O4 (LMO) Epifilms Grown on Various Single-Crystal Substrates epitaxial relationship epifilm/substrate interface

space group



(111)LMO/(111)Au (110)LMO/(110)MgO (110)LMO/(110)SrTiO3

Fd3m/Fm3m Fd3m/Fm3m Fd3m/Pm3m

orientation (cross-sectional)

parallel planes

misfit, f (%)

[112̅]LMO /[112̅0]Al2O3 [11̅0]LMO/[11̅00]Al2O3 [110]LMO/[110]Au [001]LMO/[001]MgO [001]LMO/[001]SrTiO3

(11̅0)LMO/(11̅00)Al2O3 (001)LMO/(112̅3̅)Al2O3 (110)LMO/(101̅)Au (110̅ )LMO/(110̅ )MgO (11̅0)LMO/(11̅0)SrTiO3

5.9 −1.1 1.0 −2.1 5.5

Figure 3. Cross-sectional BF-TEM images (a−d) and corresponding SAED patterns (e−h) of LiMn2O4 thin films grown on (a and e) MgO; (b and f) SrTiO3; (c and g) Al2O3; and (d and h) Au/Al2O3 single-crystal substrates at 750 °C.

around 5% is necessary to form epitaxial films in cubic/cubic heterointerface systems. In contrast, for the cubic/rhombohedral LiMn2O4/Al2O3 system, f varies in the two directions perpendicular to the interface plane because of the different crystal symmetries of the two materials., viz. cubic for LiMn2O4 and rhombohedral for Al2O3, giving f = 5.9% for (11̅0)LiMn2O4//(11̅00)Al2O3 planes,

MgO (110), aligned grains were formed even for a relatively large calculated misfit parameter of f = −2.1%, as evidenced by Figure 3a and e. At the interface, misfit dislocations are expected to form to compensate for the larger mismatch strain between LiMn2O4 and MgO.22 In the case of the SrTiO3 substrate, the ring pattern of LMO indicates that the LiMn2O4 film is polycrystalline, and a well-aligned film could not be produced, as seen from Figure 3b and f. The misfit parameter in this case is 5.5%, so that formation of an incoherent interface during CSD is not unreasonable. In contrast, there are no reports of epitaxial LiMn2O4 films prepared by pulsed laser deposition (PLD) on SrTiO3 substrates with thicknesses 20 nm or greater, because although atomic level nucleation and layerby-layer growth enable single crystals to be grown epitaxially on the substrate surface, it is a relatively slow process.12 In the CSD method, nucleation and crystallization of large volumes of precursor molecules occur simultaneously, while in the PLD method film growth occurs atom-by-atom. Consequently, when CSD is used large LiMn2O4 nuclei are unable to fully accommodate the mismatch strain, having already formed relatively large LiMn2O4 crystallites, which impedes the formation of coherent interfaces (and inhibits their ability to become fully aligned with other nuclei) when the mismatch parameter is over 5%. Taken together, these results show that epitaxial film growth by the CSD method is highly dependent on the misfit parameter and an overall mismatch of less than

but only −1.1% for (001)LiMn2O4//(112̅3̅)Al2O3 planes. The SAED pattern in Figure 3g for LiMn2O4 on an Al2O3 (0001) substrate heat-treated at 750 °C (Figure 3c) reveals that a welloriented LiMn2O4 film was produced, with an orientation relationship corresponding to (111)LMO//(0001)Al2O3 and [112̅]LMO//[112̅0]Al2O3. The SAED results thus confirm that an epitaxial LiMn2O4 thin film was successfully grown on the Al2O3 (0001) substrate, in contrast to the cubic/cubic LMO/ STO system, which has a similarly large degree of lattice misfit in both directions. This difference in structure suggests that the small misfit parameter of −1.1% between (001)LMO// (112̅3̅)Al2O3 planes can compensate for the large misfit parameter of 5.9% of (11̅0)LMO//(11̅00)Al2O3 planes in the orthonormal direction, so that the overall volumetric strain in the LiMn2O4 crystal grown on the Al2O3 substrate is less than that grown on the SrTiO3 substrate. 19543

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Figure 4. (a) Cross-sectional HAADF-STEM image taken along the [112̅]LMO/[112̅0]Al2O3 direction to show the arrangement of Mn and Al atom columns near the LiMn2O4/Al2O3 interface. A simulated interface structure is overlaid on the right of the image for comparison, and a dislocation formed at the interface is labeled “T”. (b) Inverse fast-Fourier-transform-(FFT)-filtered HAADF image of part a using the (440)LMO and (4400)Al2O3 reflections (circled in the inset showing the corresponding FFT-filtered diffraction pattern) showing clearly the position of the interface dislocation.

In an attempt to improve the quality of LMO films further, sapphire substrates were modified by depositing a thin layer of Au on the Al2O3 (0001) surface, because there is only 1.0% misfit between the lattice parameters of LiMn2O4 and Au. The Au/Al2O3 system, however, contains an interface between a cubic and a rhombohedral crystal, so that the misfit parameter is different in orthonormal directions parallel to the interface plane, similar to the case of LMO/Al2O3. The misfit parameter is 4.8% in the case of (11̅0)Au//(11̅00)Al2O3 and 2.3% in the case of (001)Au//(112̅3̅)Al2O3. This small degree of misfit enables an epitaxial (111)Au film to be successfully grown on a (0001)Al2O3 single crystal, acting as a buffer layer between the LMO film and sapphire substrate. Figure 3d and h shows that an epitaxial LiMn2O4 film with high crystallinity was successfully grown on the Au/Al2O3 (0001) substrate. Selection of an appropriate substrate with a sufficiently small misfit parameter was thus confirmed to be an important consideration in the preparation of epitaxial thin films of LiMn2O4 using the CSD technique. STEM observations were performed to examine the structure of the film, particularly the interface region, with atomic detail. HAADF-STEM has the advantage that the image contrast is roughly proportional to the square of the atomic number, Z, enabling relatively heavy elements to be distinguished for similar atomic column densities.37,38 Figure 4a shows a crosssectional HAADF-STEM image of an LiMn2O4/Al2O3 interface, together with a simulated image generated using a model of the LMO and Al2O3 crystals meeting at (111)LMO and (0001)Al2O3 crystal layers. An FFT-filtered HAADF image revealing a dislocation at the interface is also shown in Figure 4b. The image was taken along the [112̅]LMO //[112̅0]Al2O3 direction, which is parallel to the interface plane and different atomic columns, making it suitable for HAADF-STEM imaging. The epitaxial planes in this direction also have a relatively large misfit parameter. In the HAADF-STEM micrograph of Figure 4a, Mn (111) layers in LiMn2O4 can be seen to be parallel to the Al (0001) layers in Al2O3 (and the interface plane), while Mn (11̅0) and Al (11̅00) layers are aligned perpendicular to the interface plane. A dislocation formed at the LiMn2O4/Al2O3 interface can also be discerned, which is more clearly visible in the FFT-filtered HAADF image in Figure 4b. Formation of the dislocation is most likely a consequence of the relatively large lattice mismatch (5.9%) between the (11̅0)LMO and (11̅00)Al2O3 planes at the interface. Figure 4 also reveals that the terminating

layer at the interface comprises oxygen columns, since weak contrast of an oxygen layer between the Mn and Al cation layers can be discerned. Mn and Al cation columns show some streaks in their contrast, which is attributable to the presence of neighboring oxygen columns. The distance between the terminating Mn and Al planes is 0.3 nm, which is sufficiently wide to be occupied by an oxygen layer, as shown by the overlaid simulated image. The atomic model of the interface was constructed simply by joining an O-terminated plane of LiMn2O4 directly to an Al-terminated plane of Al2O3. The supercell lattice parameters were 2.324 × 1.648 × 4.000 nm3, and simulations were performed over a range of thicknesses from 5 to 20 nm. The best fit to the experimental image in the [112̅]LMO//[112̅0]Al2O3 direction was obtained when the thickness was 10 nm. These observations indicate that an epitaxial LiMn2O4 film with a semicoherent interface was successfully grown on the single-crystal Al2O3 substrate. Figure 5 shows the XRD pattern of the LiMn2O4 thin film fabricated by spin coating on the Au-coated sapphire substrate

Figure 5. X-ray diffraction pattern of an LiMn2O4 thin film on an Au/ Al2O3 substrate fabricated by coating of a precursor solution and heat treatment at 750 °C.

after heat treatment at 750 °C. The lattice spacing of the Au (111) layer on Al2O3 (0001) is 0.235 nm, which is equal to the lattice spacing in an (unstrained) Au crystal. The Au film is strongly aligned in the [111] direction above the Al2O3 (0001) substrate. The XRD pattern also shows that the LiMn2O4 is highly oriented, with (111)LMO planes parallel to (111)Au and (0001)Al2O3 planes. The lattice spacing of (111)LMO is 0.475 nm, which is similar to, but slightly smaller than, the unstrained 19544

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substrate in response to (or in order to minimize) the misfit strain, where the structural distortion is facilitated by oxygen vacancies. At the interface layer itself lithium vacancies are also present.42 The tetragonally distorted structure was observed to extend to a distance of about 10 nm from the interface, beyond which the crystal symmetry becomes cubic. The average lattice constant of the thin film was calculated to be 0.8235 nm for the cubic structure using the d-spacing of 0.4759 nm of the (111) peak in the XRD pattern of LiMn2O4. According to the XRD results, the average (400) interplanar spacing in the LiMn2O4 film is 0.2058 nm, and the (200) interplanar spacing in the Au layer is 0.2041 nm. Thus, the misfit parameter for LiMn2O4 (400) and Au (200) prepared by CSD is reduced to an average value of 0.87%, compared to the theoretical value of 1.0% given in Table 1 for an ideal cubic structure. This result is consistent with the preferential growth of tetragonal LiMn2O4 near the interface with its small misfit value and, hence, with correspondingly small lattice strain, allowing LMO thin films with excellent epitaxy to be fabricated. Figure 7 shows the cyclic voltammetry measurements obtained for the LiMn2O4 film deposited on the Au(111)/

value obtained for LiMn2O4 powder of 0.478 nm. The calculated average lattice constant of the LiMn2O4 thin film obtained from the XRD spectrum was 0.823 nm, in good agreement with that of previously reported polycrystalline samples.39,40 Figure 6a shows the projected structure model of LiMn2O4 along the [110] direction, and the cross-sectional HAADF-

Figure 6. (a) Li, O, and Mn columns observed along the [110] zone axis of the ideal LiMn2O4 spinel structure. (b) Cross-sectional HAADF-STEM image showing the atomic structure of the LiMn2O4/Au interface in an LiMn2O4/Au/Al2O3 sample. The righthand side of part b shows a HAADF image with contrast optimized to show the Au columns. The overlaid structure model of the heavy elements (Mn = yellow; Au = blue) in the two crystals shows the positions of cation columns neighboring the interface.

STEM images in Figure 6b show the atomic structure of the LiMn2O4/Au interface in the LiMn2O4/Au/Al2O3 sample. Due to the large difference in atom number between Au (Z = 79) and Mn (Z = 25), it is very difficult to show both crystal regions simultaneously in the Z-contrast HAADF image. To compensate for this, the right-hand side of the HAADF image in Figure 6 is an image whose contrast has been optimized to show the Au atom columns. Structure models have also been added to indicate the positions of Mn, Au, and Al cation columns neighboring the interface in the HAADFSTEM image. The [110] projection of LiMn2O4 is the most suitable for observing Li, O, and Mn ions directly, because separate columns of each of these ions are aligned in this direction.28 Eight Mn columns form a diamond configuration, with two distinct types of Mn columns: the number of Mn atoms in the Mn-1 columns at the vertices of the diamond is twice that of Mn atoms in Mn-2 columns between them. In the ideal cubic structure, the ratio of the projected distance along the short diagonal, m, to the distance along the long diagonal, n, is 0.707, as shown in Figure 6a. The ratio measured from HREM images of the interior of the LMO film is 0.709, very close to the m/n ratio of the ideal cubic structure. In contrast, the m/n ratio of the Mn diamond near the LiMn2O4/Au interface when measured was found to be 0.732, which is close to that obtained for tetragonal (oxygen-deficient) LiMn2O4 spinel, namely 0.73128 (tetragonal LiMn2O4 at room temperature has lattice constants a = b = 0.5735 nm and c = 0.8646 nm,41 corresponding to pseudocubic lattice parameters ac = bc = 0.8111 nm and cc = 0.8646 nm). In this case, the lattice mismatch parameter for tetragonal LiMn2O4−δ and Au is only 0.7%, which is smaller than that for cubic LiMn2O4 on Au. Tetragonal LiMn2O4‑δ is thus preferentially grown above the Au

Figure 7. Cyclic voltammogram of a LiMn2O4 epitaxial cathode film deposited on Au-coated sapphire with lithium metal as the anode and a mixture of PC and 1 M LiClO4 as the electrolytic solution. The scan rate was 0.1 mV s−1.

Al2O3 (0001). Delithiation peaks are observed at 4.03 V (A1) and 4.17 V (A2) during charging, and two lithiation peaks at 3.99 V (C1) and 4.09 V (C2) during discharge, confirming the reversibility of the Li (de)intercalation process. The two peaks are characteristic of the spinel structure and indicate that the lithium ions are extracted and inserted between LiMn2O4 and λ-MnO2 in a two-step process, consistent with previous reports.43 From Figure 7, peak separation values ΔP1 (the difference between the first anodic (A1) and cathodic (C1) peak potentials, A1 − C1) and ΔP2 (the difference between the second anodic (A2) and cathodic (C2) peak potentials, A2 − C2) were determined to be 40 mV and 75 mV, respectively. These values are smallerer than those typically reported for LiMn2O4 materials with micron-sized grains.44 Interestingly, ΔP1 is close to that reported for nanosized particles, while ΔP2 is between that of nano- and microsized particles, which may be related to the two-dimensional arrangement of crystals in the epitaxial thin film. 19545

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The Journal of Physical Chemistry C The peaks in Figure 7 are also very sharp and compare well with those reported for nanoparticle LiMn2O4 powder.44 Sharp peaks in the CV curve imply that the electrochemical reaction is completed in a very short time interval.43 The epitaxial LiMn2O4 film thus appears to support rapid lithium (de)intercalation compared with bulk polycrystalline LiMn2O4 and, thus, should be able to sustain high cycling rates, making it an attractive candidate for use as the cathode in all-solid-state lithium-ion batteries.


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CONCLUSIONS LiMn2O4 thin films formed on single-crystal cubic MgO, SrTiO3, and Au/Al2O3 substrates, as well as rhombohedral Al 2 O3 substrates, were prepared by chemical solution deposition, and the influence of the lattice misfit between the component crystals on the degree and quality of the epitaxy was compared. For the cubic/rhombohedral LiMn 2O 4/Al 2O3 system, epitaxial film growth is found to be facilitated by the formation of interface dislocations to compensate for the relative large lattice misfit along the [112̅]LMO//[112̅0]Al2O3 direction and the small misfit along a perpendicular direction, resulting in a semicoherent interface. For the cubic/cubic systems, when the misfit parameter (f) is less than 5%, epitaxial LiMn2O4 films were successfully prepared. The system LMO/ MgO, which has an f value larger than 2%, was found to form semicoherent interfaces, where a certain density of dislocations are formed to release the strain resulting from large lattice mismatch. The highest quality film with a fully coherent interface and excellent epitaxy was obtained by coating an Al2O3 substrate with Au, as the (111)LMO and (111)Au interface has the smallest misfit value of about 1.0% when cubic symmetry is assumed. Detailed observation by STEM revealed that in the interface region this lattice strain is further minimized by LiMn2O4 taking on tetragonal symmetry, facilitated by formation of oxygen vacancies, with the mismatch parameter between tetragonal LiMn2O4 and Au being smaller than that of cubic LiMn2O4. The ability of LiMn2O4 to absorb this strain through oxygen vacancy formation and increased tetragonality allows a highly coherent interface and thin film with excellent epitaxy to be formed. Cyclic voltammetry measurements of the epitaxial LiMn2O4 film on the Au/Al2O3 substrate revealed characteristic charge and discharge peaks that were sharp and with small peak separation under anodic and cathodic half-cycles. These results confirm that high-quality epifilms of LiMn2O4 can be prepared which promise to provide excellent electrochemical properties when used as the cathode in thin-film lithium-ion batteries. ASSOCIATED CONTENT

S Supporting Information *

Structure model of the LiMn2O4/Al2O3 interface used for HAADF-STEM image simulations. This material is available free of charge via the Internet at http://pubs.acs.org.


The authors thank Prof. Y. Ikuhara of the University of Tokyo and JFCC, and Dr T. Kimura of JFCC, for helpful discussions.



Corresponding Author

*E-mail: [email protected] Phone: +81-52-871-3500. Notes

The authors declare no competing financial interest. 19546

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The Journal of Physical Chemistry C


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