Nanoscale Zirconium-Abundant Surface Layers on Lithium- and

Nov 16, 2017 - Figure 3e shows the EELS spectra, obtained in HAADF–STEM mode, of an f-LMR particle at the (202̅)m surface and the bulk area. These ...
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Letter Cite This: Nano Lett. XXXX, XXX, XXX-XXX

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Nanoscale Zirconium-Abundant Surface Layers on Lithium- and Manganese-Rich Layered Oxides for High-Rate Lithium-Ion Batteries Juhyeon Ahn,†,‡ Jong Hak Kim,‡ Byung Won Cho,† Kyung Yoon Chung,† Sangryun Kim,§ Jang Wook Choi,*,∥ and Si Hyoung Oh*,† †

Center for Energy Convergence Research, Korea Institute of Science and Technology (KIST), Seoul 02792, Republic of Korea Department of Chemical and Biomolecular Engineering, Yonsei University, Seoul 03722, Republic of Korea § Institute for Materials Research, Tohoku University, Katahira 2-1-1, Aoba-ku, Sendai 980-8577, Japan ∥ School of Chemical and Biological Engineering and Institute of Chemical Processes, Seoul National University, Seoul 08826, Republic of Korea ‡

S Supporting Information *

ABSTRACT: Battery performance, such as the rate capability and cycle stability of lithium transition metal oxides, is strongly correlated with the surface properties of active particles. For lithium-rich layered oxides, transition metal segregation in the initial state and migration upon cycling leads to a significant structural rearrangement, which eventually degrades the electrode performance. Here, we show that a fine-tuning of surface chemistry on the particular crystal facet can facilitate ionic diffusion and thus improve the rate capability dramatically, delivering a specific capacity of ∼110 mAh g−1 at 30C. This high rate performance is realized by creating a nanoscale zirconium-abundant rock-salt-like surface phase epitaxially grown on the layered bulk. This surface layer is spontaneously formed on the Li+-diffusive crystallographic facets during the synthesis and is also durable upon electrochemical cycling. As a result, Li-ions can move rapidly through this nanoscale surface layer over hundreds of cycles. This study provides a promising new strategy for designing and preparing a high-performance lithium-rich layered oxide cathode material. KEYWORDS: Li- and Mn-rich layered oxides, transition metal segregations, Zr-abundant surface layers, rate capabilities, nanoscale, crystallographic facets

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mation of the surface structure is closely associated with the segregation of TM atoms like Ni and Co at the outermost surface of the LMR particle during the synthesis.27,28 The local concentration of such TM atoms also varies along the LMR surface; a significant enrichment of Ni or Co atoms are observed at particular surface orientations.9,24,26,29 More specifically, Yan et al.9,26 demonstrated that the segregation of Ni and Co in LMR particles is strongly selective for particular crystallographic facets that are open to Li diffusion channels in the bulk particles. These TM-segregated surface structures often act as ionically insulating layers, triggering a large impedance buildup.17,18,30 As a result, LMR undergoes severe capacity fading at high current rates.18,31 To alleviate these surface reconstructions and transition metal enrichments in the surface region, various syntheses and surface modifications were applied up to now for the elaborate control of the surface structures. Zhang et al. prepared a

hile lithium transition metal oxides are dominantly adopted for current commercial lithium-ion batteries (LIBs), lithium- and manganese-rich nickel-cobalt-manganese oxides (LMRs) with the chemical formula Li1+x(Ni, Mn, Co)1−xO2 are receiving discernible attention because of their superior specific capacities, often exceeding 250 mAh g−1.1−5 When LMRs are charged to high voltages (>4.5 V versus Li+/ Li), Li-ion removal and oxygen release occur simultaneously, accompanied by a gradual phase transition to a spinel or rocksalt structures or both, primarily due to the transition metal (TM) migration into the Li layers.6−8 This undesirable structural transformation proceeds from the surface of the particles9−18 and undermines the ion diffusion characteristics significantly, resulting in poor rate capability as well as voltage and capacity fading over prolonged cycles.4,7,17−19 On that account, the use of LMRs for practical applications has been limited so far.20 Besides the TM migration upon electrochemical cycling, the surface structure in the pristine material has been also recognized as one of the major parameters that determine the electrochemical performances of LMR.9,21−27 The for© XXXX American Chemical Society

Received: September 27, 2017 Revised: November 7, 2017

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DOI: 10.1021/acs.nanolett.7b04158 Nano Lett. XXXX, XXX, XXX−XXX

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Figure 1. Characterization of the crystal habits of f-LMR and s-LMR. (a,d) SEM images of f- and s-LMR powders. (b,e) Bright-field TEM images of f- and s-LMR particles. An f-LMR particle is bounded by the (001)m, (200)m, and (202̅)m surface planes. For s-LMR, no particularly developed facets are observed except the (001)m planes. (c,f) Selected-area electron diffraction (SAED) patterns of panels b and e, respectively, and can be indexed to [010]m zone of monoclinic lattice with space group C2/m.

segregation-free LMR with a homogeneous bulk-to-surface structure through a hydrothermal method.31 Numerous approaches were also reported to stabilize the surface structure of LMR including the surface doping32,33 and surface coating with metal oxides,2,34,35 metal fluorides,36,37 or graphene oxides.38 To endow the surface layer with ion transport properties, Li-ion conducting metal oxides (e.g., Li2XO3, X = Zr, Ti, or Si) were introduced on the LMR particle surface.39−42 These were typically accomplished through well-established wet and dry coating processes such as sol−gel and spray drying. Recent studies by Xie et al. achieved a uniform surface coating with Li-ion conductive metal oxides and fluorides (e.g., LiAlO2 and LiAlF4) on the layered oxide electrodes using atomic layer deposition.43,44 Although these attempts employing a surface modification of the as-synthesized materials were effective to some extent, more integrated approaches controlling both the nature and the interface of the surface chemistry, e.g., a segregation layer, could bring a drastic improvement in the ion transport behavior through the surface while largely maintaining the host structure. Herein, we developed Li- and Mn-rich transition metal oxide particles with a few nanometer-thick zirconium-abundant surface layer, which was spontaneously and epitaxially grown on Li+-diffusive (202̅)m facets via a one-pot synthesis and also durable upon an electrochemical cycling. Our strategy is that on the introduction of zirconium, a rock-salt-like oxygen-deficient surface phase is concurrently created on the particular surface planes of the LMR, while suppressing the unwanted metal segregation in the initial state. The atomic-scale characteristics of this surface layer were probed by advanced analytical techniques, including high-resolution aberration-corrected

scanning transmission electron microscopy (STEM) with a high-angle annular dark-field (HAADF) detecting mode, electron energy loss spectroscopy (EELS), and energydispersive X-ray spectrometry (EDS). The electrochemical studies showed that LMR with this plane-selective nanoscale surface layer exhibited an exceptional high-rate performance and long-term cycling stability. Results and Discussion. In this work, Li- and Mn-rich layered oxide particles (Li1.2Ni0.13Co0.13Mn0.54O2, LMR) with nanoscale surface layers were prepared by a one-pot synthesis using zirconium acetate hydroxide as a complexing agent (see the Methods section for synthesis details).45 The morphologies of the as-prepared LMR particles depend critically on the presence of potassium nitrate in the precursor solution. The addition of potassium nitrate yields polyhedra with welldeveloped crystallographic facets (denoted as f-LMR), as seen in the scanning electron microscopy (SEM) and transmission electron microscopy (TEM) images shown in Figure 1a−c. Without potassium nitrate, however, the primary LMR particles are rather spherica,l with diameters of 200−300 nm (denoted as s-LMR), as shown in Figure 1d−f. The crystal structure of Li- and Mn-rich transition metal oxides is somewhat elusive.46,47 Their structures are described as xLi2MnO3·(1-x)LiMO2 (0 < x < 1, M = Ni, Co, and Mn), consisting of either a composite of a monoclinic (m) Li2MnO3like phase with C2/m symmetry and a rhombohedral (r) LiMO2 phase with R3̅m symmetry or a solid-solution of the monoclinic phase with C2/m symmetry. In these two structural models, the local atomic distributions, such as the partial presence of lithium in the transition metal layer, might be slightly different. However, both descriptions adopt the layered structures in B

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Figure 2. Atomic-scale observation of surface and bulk structure of f-LMR. (a) Z-contrast image for the bulk-to-surface region near (202̅)m facet of fLMR projected along the [010]m zone axis. The inset shows where the HAADF-STEM image was taken. (b) Corresponding HAADF-STEM signal profiles at the surface region (labeled as region I in panel a). (c) Corresponding HAADF-STEM signal profiles at the bulk region (labeled as region II in panel a).

constant ratio is greater than 4.899, indicative of ideal cubic close packing54,55 (Table S1 and Figure S3). In addition, the intensity ratio of the (003)r peak to the (104)r peak (I(003)/ I(104)) is greater than 1.2 for both f-LMR and s-LMR, indicating a very low degree of cation mixing in the bulk.54−56 However, the structure and atomic distribution at the outermost surface and subsurface region may be vastly different from those of the bulk, and, in this study, we directly investigated these properties microscopically. Figure 2a is a Z-contrast image along the [010]m direction acquired from a high-resolution aberrationcorrected STEM with a HAADF detector. The contrast of the HAADF−STEM image is proportional to the atomic number (Z1.7);57,58 that is, a brighter spot indicates the presence of heavier elements. Thus, the lithium and oxygen atoms are relatively invisible because of their low atomic numbers. In contrast, the 3d TM atoms and 4d TM Zr atoms can be clearly seen as bright spots. In pristine s-LMR, the TM and Li layers are consistently distributed from the bulk to the surface of a particle without significant cation disorder in the octahedral sites of the respective layers (Figure S4). Contrary to the homogeneous bulk-to-surface structure in s-LMR, in f-LMR, heavy metal atoms (i.e., Ni, Co, Mn, and Zr) are seen in the Li slabs within 1−2 nm of the (202̅)m surface planes (Figures 2a and S5 and Table S2). This cation-disordered surface phase was observed only in the particles with clearly developed facets. Figure 2b,c shows the intensity profiles of the HAADF images across the surface layers (region I) and the bulk phase (region II) of an f-LMR particle, respectively. For the bulk, the intensities of the Li layers between the TM layers are uniformly low, which indicates that Li and TM cations occupy their appropriate sites, respectively. The interplanar spacing of the (001)m (or (003)r) plane was measured to be ∼4.7 Å, which is in close accordance with the result determined by X-ray diffraction (d(003)r = 4.747 Å). However, at the surface, the

which Li-ions predominantly diffuse parallel to the (001)m (or (003)r) planes.48−51 The projection images for both C2/m and R3m ̅ phases are almost the same when viewed along the [010]m/[110]r direction.22,47 In fact, irrespective of the actual nature of the LMR crystal structure, it is crucial to identify the surface facets through which Li-ions can pass. Consequently, the atomic-level observations are generally projected along the [010]m direction. Figure 1b and c show the bright-field TEM image and a selected-area electron diffraction (SAED) pattern of an f-LMR particle, respectively. The exposed surface facets of the f-LMR particle are identified as the (001)m, (200)m, and (202̅)m crystal planes (see Figure S1 to locate these planes in the unit cell). Among the developed facets, the (202̅)m planes are aligned to Li-ion diffusion paths. In contrast, the s-LMR particles do not show any noticeable facet development at the surfaces aligned to the Li-ion insertion and extraction (Figure 1e,f). While the (001)m planes are formed in both f-LMR and sLMR, the development of the surface facets toward diffusion paths are dependent on synthesis conditions, which is consistent with the work of Wei et al.,52 in which the surface energy of the (001)m planes was estimated to be lower than that of the (010)m planes in Li- and Mn-rich layered oxide materials. They claimed that the (010)m planes disappear gradually during the growth process because of their high surface energies, while the (001)m planes are maintained. In our case, however, during the preparation of f-LMR, K+ ions are speculated to alter the surface energy of some planes like the (202̅)m planes and, thus, affect the growth rate of the crystal planes and the final morphologies of the particles.53 This trend was confirmed in the multiple LMR particles (Figure S2). Layered metal oxides with an α-NaFeO2 structure are conventionally described as alternating stacks of lithium, oxygen, and TM layers. The crystal structure of bulk f-LMR and s-LMR are highly cation-ordered because the c-to-a lattice C

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Figure 3. Relative distribution of chemical elements and EELS spectra for a single f-LMR particle. (a) High-resolution HAADF-STEM image along the [010]m zone axis. (b) EDS maps of individual elements including Mn, Co, Ni, and Zr. Relative atomic concentration from the surface of (c) Zrsegregated (202̅)m plane and (d) (200)m plane to the bulk of a particle of a pristine f-LMR particle in panel a. (e) Comparison of EELS spectrum profiles at the (202̅)m surface (red) and in the bulk region (blue) of f-LMR. STEM-EELS spectra for O K-edge, Mn L3,2-edge, Co L3,2-edge, and Ni L3-edge at the Zr-segregated surface and the bulk region of a pristine f-LMR are obtained.

disordered rock-salt-like phase has been epitaxially grown on the (202)̅ m facets of the layered bulk structure.60,61 Figure 3 shows the elemental distributions (Ni, Mn, Co, and Zr) obtained by EDS mapping and local oxidation states at surface and in the bulk of a single f-LMR particle. The HAADF image in Figure 3a illustrates that the f-LMR particles hold (001)m, (200)m, and (202̅)m surface planes when viewed from the [010]m direction. As clearly displayed in Figure 3b, Zr atoms segregate mostly along the (202̅)m facets. In contrast to the f-LMR particles, Zr segregation is not pronounced following the surfaces of s-LMR particles. Rather, the Zr atoms are homogeneously distributed throughout the entire particle (Figure S7). Figure 3c shows the relative distributions of chemical elements within a few nanometers of the surface toward the bulk in the [202̅]m direction, as indicated by area marked in red in Figure 3a. Within an extremely thin surface layer (approximately ∼5 nm thick), the fractional amount of Zr sharply increased to ca. 15 atomic percent (at. %) toward the end of the particle, reconfirming the aggregation of Zr atoms in

intensities of Li slabs abruptly increase, implying that Li-ions no longer solely exist in the slabs. We speculated that a significant number of transition metals had been introduced to the Li layers; consequently, the interplanar spacing of the two adjacent TM layers (2.4 Å) becomes approximately half of that from the bulk area (4.7 Å). At the outermost surface region, the contrast intensity of Li slabs and TM slabs are similar, which implies that cations are randomly distributed in either Li or TM sites. To verify the structural inhomogeneity near the surface region, we attained the fast Fourier transform (FFT) patterns of the surface (region III) and bulk area (region IV) images, as shown in panels b and c of Figure S6, respectively. While the FFT result of the bulk area is consistent with that of the layered structure, the FFT pattern of the surface layer is indicative of a cubic rock-salt structure,59 similar to previous reports10,11,14,60 concerning TM segregation along the surfaces of LMR particles. Combined with the results of Zcontrast STEM analysis, this result indicates that a cationD

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Figure 4. Electrochemical performances of f-LMR and s-LMR. (a) The first charge and discharge profiles of f- and s-LMR at 0.1C rate, where a 1C rate corresponds to the current density of 200 mA g−1. (b) Rate capabilities of f- and s-LMR electrodes at various current rates of 0.1, 0.5, 1, 2, 5, 10, 20, and 30C rates after the activation at the 0.1C rate for one cycle. (c) The charge and discharge voltage profiles of f- and s-LMR at 0.1, 1, and 10C rates. (d) Cycling stabilities of f- and s-LMR at a 30C rate for 300 cycles after the activation at C/5 rate for two cycles. All electrochemical measurements are carried out in a galvanostatic mode at 25 °C.

electrochemically detrimental because the resultant insulating rock-salt-like NiOx phase impedes lithium ion diffusion by blocking its pathways.30,38 Our study clearly shows that, in fLMR, the Zr-abundant phase effectively inhibits the segregation of 3d TM atoms along the (202̅)m planes and is therefore expected to have a positive impact on the electrochemical properties of f-LMR. Figure 3e shows the EELS spectra, obtained in HAADF− STEM mode, of an f-LMR particle at the (202̅)m surface and the bulk area. These EELS spectra elucidate the local environments of oxygen and TMs. All spectra were normalized with respect to the intensity of the main peak. For the oxygen K-edge, the pre-edge peak at around 527−530 eV arises from the electron transition from O 1s core state to the unoccupied O 2p states hybridized with TM 3d states.63 It is known that the reduced intensity of the O pre-edge is mainly attributed to the formation of oxygen vacancies.64,65 Because the intensity of the O pre-edge in the surface region is much lower than that from the bulk area, we concluded that the oxygen vacancies exist in the surface region. The EELS analysis unveiled the local electronic states of Mn, Co, and Ni as well. The L-edge spectrum for the 3d TM consists of two white lines, L3 and L2, arising from the transitions from the 2p3/2 and 2p1/2 core states to the 3d unoccupied states.60,63,66 The valence states of the TMs can be evaluated by referring to either the peak shifts or the intensity ratios of the L3 and L2 white lines.15,18,34 The L3 peaks for all

the surface region. Note that Zr content among TMs in LMR is less than 1 at. % (Table S3). STEM−EDS analysis also supports that the Zr atoms are quantitatively more concentrated in the surface region along the (202̅)m planes than in the bulk (Figure S5). The facet-specific Zr occupation on the surface sites might be explained by the fact that the formation of Zr-abundant phase lowers the surface energy of the (202)̅ m planes.24,27,52 The presence of Zr also prevents the segregation of 3d TM atoms on the (202̅)m facets during material synthesis (Figure 3c), as the given high-energy facets would have otherwise been stabilized by other TM atoms, in the absence of Zr.62 In the (200)m facet devoid of Zr segregation, however, high concentrations of Ni and Co are observed (Figure 3d). TM segregation and the resulting crystal structure at the pristine LMR surface layer is known to have a significant impact on the electrochemical behavior.10,32,38 Previous reports demonstrated that TM segregation in the pristine particles depends on both the chemical composition and the synthetic method used. Zhang’s group concluded that, on the surface of Li1.2Ni0.2Mn0.6O2 synthesized by co-precipitation, substantial Ni segregation occurs along the (200)m planes.9 More recently, they additionally found that Ni and Co segregate along the (200)m and (202̅)m planes of Li1.2Ni0.13Co0.13Mn0.54O2 particles prepared by a molten salt method.26 These studies coherently indicate that Ni and Co are preferentially located along the certain facets. However, the segregation of 3d TM atoms is E

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Figure 5. STEM-EDS maps of Zr of an f-LMR particle upon electrochemical cycling. (a) A particle at the pristine state, (b) a particle after initial charge, and (c) a particle after 2 cycles at a 0.1C rate and then 20 cycles at a 1C rate.

shifts for f-LMR are much smaller from their equilibrium positions compared with those of s-LMR, which is another signature of the lower electrochemical impedance of f-LMR (Figure S9).64 Figure 4d shows the cycling performance of the f-LMR cathode when evaluated at a high current rate of 30C. Notably, the f-LMR cathode delivered high discharge capacity (∼110 mAh g−1 at 6 A g−1) and maintained 85% of its maximum capacity, even after 300 cycles. This result demonstrates that the Li- and Mn-rich oxide cathodes can sustain their operation over a great number of cycles by engaging durable surface phases on the Li+-diffusive crystallographic facets of the host particles. Furthermore, we evaluated the stability of this nanoscale surface layer upon the electrochemical cycling. Figure 5a−c show the STEM-EDS mapping images of f-LMR at the pristine state, after initially charging to 4.8 V, and after 20 cycles at 1C, respectively. This surface layer across (202̅)m facet was well-preserved after these tests. It is known that in the case of Li- and Mn-rich layered transition metal oxides, TMs are often segregated preferentially along particular surface facets. Based on this observation, we developed a nanoscale Zr-abundant surface layer on the Li+diffusive facets by employing a one-pot synthetic method including the Zr-containing precursor. This nanometer-scale surface layers along (202̅)m planes are aligned with Li-ion diffusion pathways and also suppress the undesirable segregation of TMs. As a consequence, electrochemical performance of LMR was improved dramatically, especially toward supporting high C-rate operation (i.e., the 30C rate). The series of results imply that the Zr-containing surface layer perform a dual function of mitigating TM aggregation and supporting Li-ion diffusion through well-aligned channels orientations. The observation of the fast Li-ion transport through the surface layer of f-LMR with a rock-salt-like structure can be understood in the same line as recent rationale of Ceder et al. that Li-ion conduction in Li-/Mn-rich transition metal oxides with disordered rock-salt structures follows the percolation routes.70,71 Conclusions. In summary, we have developed the Li- and Mn-rich layered oxides with a nanoscale zirconium-segregated rock-salt-like surface phase that is arranged along the Li+diffusive facets of the bulk particles and also durable upon electrochemical cycling. These features enable an ultrahigh rate capability and stable performance in the long-term electrochemical cycling. We believe that our findings in this work reveal the importance of controlling the subtle surface

TMs are shifted to lower energies, indicating a decrease in the oxidation state at the surface compared to that in the bulk. Specifically, the Mn L3-to-L2 intensity ratio is approximately 2.7 at the surface, which is conspicuously higher than that in the bulk (1.7). This indicates that the Mn atoms are mostly divalent at the surface but tetravalent in the bulk, which agrees well with previous reports.10,18,67,68 The oxidation states of Co and Ni are also found to be divalent.25,63,68,69 These difference in the oxidation states of 3d TM at (202̅)m surface from the bulk area may account for the formation of oxygen vacancies at the surface region.25,65 To investigate the impact of the nanoscale surface layer on the electrochemical performance of Li- and Mn-rich transition metal oxides, electrochemical cells containing f-LMR cathodes were evaluated at various current rates and over prolonged cycles. The performance of an s-LMR cathode bearing neither a unique surface layer nor TM segregation (Figure S7) was also examined for comparison. As shown in Figure 4a, the initial charge voltage profiles of f-LMR and s-LMR at a C/10 rate (1C = 200 mA g−1) are almost identical, indicating that the activation of the Li2MnO3 phase occurs similarly for both materials. Figure 4b illustrates the rate performance of both active materials at various current rates from C/10 to 30C. At a low current rate (C/10), both materials show consistently high discharge capacities of around 245 mAh g−1. However, at current rates higher than C/2, the discharge capacities of fLMR are significantly larger than those of s-LMR (Figure S8); for example, f-LMR has a capacity of 130 mAh g−1 at a 10C rate, which is approximately 65% of the capacity at a 1C rate, while s-LMR exhibits only 105 mAh g−1 at the same 10C, which is only 59% of the capacity at a 1C rate. This is one of the highest performances reported to date in terms of the rate capability of the LMR cathodes (see Table S4 for a comparison with literature values). Once again, the high rate capability of fLMR is ascribed to the nanoscale surface phase, which facilitates facile Li-ion diffusion to the bulk and suppresses the segregation of 3d TM atoms on the surface region. On the whole, the Zr-containing surface layer significantly reduces the charge-transfer impedance on the solid−electrolyte interface, which will be discussed in detail below. The charge and discharge voltage profiles with increasing current rates are shown in Figure 4c. At the beginning of discharge, the voltage drops at high current rates for f-LMR are generally much smaller as compared with those for s-LMR. The corresponding differential capacity versus voltage curves (dQ dV−1) also clearly demonstrate that the cathodic or anodic peak F

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(PANAX ETEC). The electrochemical measurements of the cycling properties were performed in constant current modes using a battery and cell test equipment (Series 4000, Maccor Inc.) between 2.0 and 4.8 V versus Li/Li+ at 25 °C. The high rate performance at 30C was measured after two cycles of activation at a C/5 rate. The 1C current rate was set to 200 mA g−1.

properties of lithium transition metal layered oxides. Furthermore, this understanding will serve as a fundamental basis for the design of future high-performance energy-storage materials. Methods. Synthesis of LMR. To synthesize f-LMR, stoichiometric amounts of LiNO3, Ni(NO3)2·6H2O, Co(NO3)2·6H2O, and Mn(NO3)2·4H2O were dissolved into a solution containing zirconium(IV) acetate hydroxide (0.5 wt %) ((CH3CO2)xZr(OH)y, x + y ≈ 4). Potassium nitrate (0.028 M) was simultaneously added to the precursor solution. The solution was then sonicated for 90 min to obtain a clear solution. After solvent evaporation at 100 °C for 3 h, a gel-like substance was collected and was then subjected to a heat treatment at 600 °C for 5 h, followed by calcination at 950 °C for 10 h in air. The resulting Li1.2Ni0.13Co0.13Mn0.54O2 powders were gently ground using an agate mortar and pestle. To prepare s-LMR, the same procedure was followed, but KNO3 was not added. Material Characterization. The crystallographic structure was examined by X-ray diffraction using Cu−Kα radiation (D/ MAX-2500 V, Rigaku) and the lattice parameters were calculated using the GSAS program. The morphology was observed by a field-emission SEM (FEI NOVA NanoSEM200). TEM Characterization. TEM samples were prepared by dropping a sonicated solution of LMR powder and anhydrous ethanol onto a lacey carbon 200-mesh copper grid (Electron Microscopy Sciences). High-resolution TEM and HAADF− STEM were carried out using a Cs-corrected FEI Titan 80−300 microscope operated at an accelerating voltage of 300 kV. Images were recorded by a 2000 × 2000 CCD (Gatan UltraScan 1000) cameras. The probe resolution for STEM mode was 0.14 nm at an operating accelerating voltage. Elemental analyses were carried out using an EDS spectrometer (Super-X EDS detector) attached to the TEM equipment. EELS spectra were measured using a Gatan Quantum 966 spectrometer. The energy resolution of the system was 0.9 eV. The onset energy was carefully calibrated in advance before each elemental measurement. Gatan DigitalMicrograph and TEM imaging and analysis (TIA) software were used to analyze the images and spectra acquired with the FEI transmission electron microscope. Fast EDS mapping was conducted using an FEI Talos F200X with an accelerating voltage of 200 kV. Data analyses including spectrum acquisition, image capture, elemental mapping, and line scans were processed by the Quantax system from the Bruker Esprit program. Electrochemical Measurements. A composite of 90 wt % as-prepared power samples, 6 wt % acetylene carbon black (Denka Black, Denki Kagaku), and 4 wt % polyvinylidene fluoride (PVDF) dissolved in N-methyl-2-pyrrolidone (NMP) was mixed in a planetary ball mill (Pulverisette 23 Mini-Mill, Frisch Gmbh) at 2400 oscillations per minute for 30 min. Electrodes were prepared by casting the slurry onto an aluminum current collector. After drying at 80 °C, the electrodes were pressed to 70% of the original thickness. Next, they were dried at 80 °C in a vacuum oven for 24 h to remove residual moisture. The half-cells were assembled along with a metallic lithium anode using 2032-type coin cells in a dry room (dew point as low as −100 °C). The average loading of the active material was approximately 2.2−2.6 mg cm−2. Polypropylene (Celgard 2500) was used as a separator. The electrolyte was a 1.0 M lithium hexafluorophosphate (LiPF6) solution in a 1:1:1 (v/v %) mixture of ethylene carbonate (EC), diethyl carbonate (DEC), and dimethyl carbonate (DMC)

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.nanolett.7b04158. Crystal structure and additional results from XRD, ICP, HRTEM and STEM-HAADF analyses; EDS mapping; and electrochemical characterization for the comparison between different synthetic conditions (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. ORCID

Juhyeon Ahn: 0000-0002-1701-3372 Jong Hak Kim: 0000-0002-5858-1747 Kyung Yoon Chung: 0000-0002-1273-746X Sangryun Kim: 0000-0001-8617-3022 Jang Wook Choi: 0000-0001-8783-0901 Si Hyoung Oh: 0000-0002-7063-9235 Funding

This work was supported by the National Research Foundation of Korea (grant no. NRF-2011-C1AAA001-0030538) and KIST Institutional Program (grant no. 2E27062). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors gratefully thank Ms. M. K. Cho in the Advanced Analysis Center at the Korea Institute of Science and Technology for contributions to the microstructure characterization using TEM. We also thank Dr. S. Hwang at the Korea Institute of Science and Technology for valuable discussion on TEM analysis. We greatly appreciate Prof. J. S. Kim in Dong-A University for the general comments on this subject.

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REFERENCES

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DOI: 10.1021/acs.nanolett.7b04158 Nano Lett. XXXX, XXX, XXX−XXX