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Hierarchically Assembled 2D Nanoplates and 0D Nanoparticles of Lithium-Rich Layered Lithium Manganates Applicable to Lithium Ion Batteries Ja Yeon Baek,† Hyung-Wook Ha,† In-Young Kim, and Seong-Ju Hwang* Center of Intelligent Nano-Bio Materials (CINBM), Department of Chemistry and Nano Sciences, Ewha Womans UniVersity, Seoul 120-750, Korea ReceiVed: May 1, 2009; ReVised Manuscript ReceiVed: August 20, 2009
The porous hierarchical assembly of lithium-rich Li1+xMnO3-δ 2D nanoplates as well as isolated 0D nanocrystalline homologues has been synthesized via lithiation reactions of nanostructured manganese oxides under hydrothermal conditions. According to powder X-ray diffraction and electron microscopy, a hydrothermal LiOH treatment for nanostructured δ-MnO2 precursor produces a lithium-rich Li1+xMnO3-δ phase with the nanoworm-like hierarchically assembled 2D nanoplate morphology. After the lithiation reaction under identical conditions, the 1D nanowires of the R-MnO2 precursor are transformed into the 0D nanoparticles of the Li1+xMnO3-δ phase. The Mn K-edge X-ray absorption spectroscopic analysis for the lithiated materials clearly demonstrated that tetravalent manganese ions are stabilized in octahedral sites of a Li2MnO3-type layered structure composed of edge-shared MnO6/LiO6 octahedra. From electrochemical measurements, it was found that the lithiated Li1+xMnO3-δ nanostructured materials show much superior electrode performance over the precursor manganese oxides and bulk lithium-rich manganate. The powder X-ray diffraction analyses for the electrochemically cycled derivatives clearly demonstrated that the improvement of electrode performance after lithiation can be attributed to the phase transformation to the Li-rich Li1+xMnO3-δ phase with high structural stability. On the basis of the present experimental findings, we are able to conclude that the present phase transformation route provides a new method not only to synthesize nanostructured lithium-rich manganese oxides with controllable dimensionality and morphology but also to improve the electrode performance of nanostructured manganese oxides. Introduction Lithium manganese oxides have attracted intense research interest as alternative electrode materials for lithium rechargeable batteries because of their economic and ecological merits.1-4 Many structure types of lithium manganese oxides have been developed as lithium intercalation electrodes, i.e., spinel LiMn2O4, orthorhombic LiMnO2, lithium-rich monoclinic Li2MnO3, and so on.5-13 Most of the research ever reported have concentrated on well-crystalline lithium manganate microcrystals prepared by heat-treatment at elevated temperatures.1,2,5,6 Recently, nanocrystalline lithium manganese oxides received special attention as new efficient electrode materials applicable to large-scale battery application.14-19 The small particle size of the nanocrystals makes the diffusion length of lithium ions shorter during the electrochemical charge-discharge process, resulting in the improvement of electrode performance. There are considerable numbers of works regarding the synthesis and electrode application of low-dimensional nanostructured manganese oxides.14,15,20,21 For instance, we have reported a series of works about the composition- and structure-tunable synthesis of nanostructured manganese oxides and their application in lithium secondary batteries.22-25 In contrast to binary manganese oxides, only a few works were carried out for lithium-containing ternary manganese oxides with low dimensional nanostructure.21,26 This would be due to difficulty in incorporating lithium ions in the crystallization step of manganate nanostructure. Previously, we were successful in preparing Lix[Mn1-yNiy]O2 nanowires * To whom correspondences should be addressed. Tel: +82-2-3277-4370. Fax: +82-2-3277-3419. E-mail:
[email protected]. † These two authors contributed equally to this work.
through the chemical oxidation of a layered lithium nickel manganate precursor.22 However, since the oxidation of the precursor material led to a partial loss of lithium ions, this oxidation route is not suitable for synthesizing other types of nanostructured lithium manganate such as the Li-rich Li2MnO3 phase. In order to circumvent this difficulty, we have suggested the lithiation of nanostructured manganese oxides as an alternative route to nanostructured lithium manganate. Here, we report the morphology-controllable synthesis of nanostructured lithium-rich layered lithium manganates via the lithiation of manganese oxide precursors under hydrothermal conditions. The crystal and electronic structures and crystal morphology of the resulting nanomaterials were systematically investigated, along with their electrode performance and structural variations after electrochemical cycling. The adoption of harsh lithiation conditions with the hydrothermal vessel results in the formation of lithium-rich Li1+xMnO3-δ phases, instead of spinel LiMn2O4 with lower lithium content. To the best of our knowledge, the porous hierarchical assembly of 2D nanoplates has never been reported not only for lithium-rich lithium manganate but also for other types of lithium manganates. Experimental Section Preparation. The precursor δ-MnO2 with hierarchically assembled 2D nanoplate morphology was prepared by hydrothermal treatment for an aqueous solution containing 0.01 mol KMnO4, 0.001 mol (NH4)2S2O8, and 0.001 mol MnSO4 at 100 °C for 72 h. The precursor of the R-MnO2 1D nanowire was prepared by the hydrothermal treatment for an aqueous solution containing 0.01 mol KMnO4 and 0.01 mol (NH4)2S2O8 at 140
10.1021/jp904072r CCC: $40.75 2009 American Chemical Society Published on Web 09/11/2009
Hierarchically Assembled 2D Nanoplates °C for 24 h. The nanostructured compounds of the Li-rich Li1+xMnO3-δ phase were synthesized by the hydrothermal treatment for the aqueous mixtures of excess LiOH and nanostructured manganate precursors with δ-MnO2 or R-MnO2 structure at 160 °C for 24 h. After each hydrothermal reaction, all of the products were washed thoroughly with distilled water and dried at 50 °C in air. Characterization. The crystal structures of all of the present materials were identified by powder X-ray diffraction (XRD) using Ni-filtered Cu KR radiation with a graphite diffracted beam monochromator. The crystal morphologies of these nanostructured materials were probed by field emission-scanning electron microscopy (FE-SEM, Jeol JSM-6700F) and high resolution-transmission electron microscopy (HR-TEM, PhilipsCM200, 200 kV). The ratio of Li/Mn was analyzed with an inductively coupled plasma spectrometer (ICP, Optima 4300DV). The N2 adsorption-desorption isotherms of the lithium manganates were measured with an ASAP 2020 machine. X-ray absorption spectroscopy (XAS) experiments were carried out at the Mn K-edge using an extended X-ray absorption fine structure (EXAFS) facility installed at the Beamline 7C at the Pohang Accelerator Laboratory (PAL) in Pohang, Korea. The XAS data were collected at room temperature in a transmission mode using gas-ionization detectors. All of the present spectra were calibrated by measuring the spectrum of Mn metal foil simultaneously. The data analysis for the experimental spectra was done by the standard procedure reported previously.27 In the course of the EXAFS fitting analysis, the coordination number (CN) was fixed to the crystallographic values, while the amplitude reduction factor (S02) was allowed to vary. All of the bond distances (R), Debye-Waller factors (σ2), and energy shifts (∆E) were set as variables, under the constraint that the energy shifts were kept at the same value for adjacent (Mn-Mn) shells in the R-MnO2 material. Such constraints can be rationalized from the fact that the adjacent shells consisting of the same types of atoms at very close distances would possess nearly the same degree of energy shift. The electrochemical measurements were performed with the button-type cell of Li/1 M LiPF6 in ethylene carbonate/diethyl carbonate (50:50 v/v)/ active material, which was assembled in a drybox. The composite cathode was prepared by thoroughly mixing the active material (70%) with 20% of acetylene black and 10% of PTFE (polytetrafluoroethylene). All of the experiments were carried out in a galvanostatic mode with a Maccor multichannel galvanostat/potentiostat in the voltage range of 2.0-4.8 V at a constant current density of 20 mA/g. To investigate the redox behaviors of the present materials, cyclic voltammetry (CV) was carried out using the above-mentioned button-type cells in the voltage range of 2.0-4.8 V with a scan rate of 1 mV/s. Results and Discussion Powder XRD, FE-SEM, and HR-TEM/SAED Analyses. The variations of the crystal structures of manganate precursors upon hydrothermal lithiation reaction were examined with powder XRD analysis. The powder XRD patterns of the precursors of δ-MnO2 and R-MnO2, and their hydrothermally lithiated derivatives are plotted in Figure 1. Upon the lithiation reaction, both the precursors of δ-MnO2 and R-MnO2 experience a distinct phase transition to lithium-rich Li2MnO3-type structure with C2/m symmetry.28 A splitting of the (018)/(110) doublet near 2θ ) 65° obviously indicates the formation of a wellordered layered structure, rather than cubic spinel LiMn2O4 structure.29 Furthermore, a broad asymmetric peak at 2θ ) ∼23° (denoted as asterisk), generally assigned as a superlattice
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Figure 1. Powder XRD patterns of (a) the precursor δ-MnO2 and (b) its hydrothermally lithiated derivative, and (c) the precursor R-MnO2 and (d) its hydrothermally lithiated derivative. The asterisk symbol represents the superlattice line of the Li-rich manganate phase.
Figure 2. FE-SEM images of (a) the precursor δ-MnO2 and (b) its hydrothermally lithiated derivative, and (c) the precursor R-MnO2 and (d) its hydrothermally lithiated derivative.
ordering of Li and Mn in Li2MnO3 phase, can be regarded as strong evidence for the formation of lithium-rich layered lithium manganates, since this peak can be observed only for the Lirich manganate phase.28,30 We have examined the influence of hydrothermal lithiation reaction on the crystal morphology of the precursor materials using FE-SEM. Figure 2 represents the FE-SEM images of the precursors and hydrothermally lithiated products. The nanoworm-like hierarchical morphology of the precursor δ-MnO2 formed by assembled 2D nanoplates is well-maintained after the hydrothermal LiOH treatment. The hydrothermal LiOH treatment leads to an increase in the thickness of component 2D nanoplates from 10 to 30 nm but does not change the diameter of the secondary hierarchical spheres (i.e., 900-1000 nm). However, the lithiation process for the 1D nanowire of precursor R-MnO2 causes a remarkable morphological variation to spherical 0D nanocrystals with the diameter of ∼50 nm. Such different evolution of precursor morphology upon lithiation can be understood in terms of the structural relationship between the precursors and the products. That is, since the layered structure of precursor δ-MnO2 is basically the same as that of the lithiated Li1+xMnO3-δ, the layered δ-MnO2 precursor can be transformed into the layered Li1+xMnO3-δ without significant change in morphology. In contrast, the R-MnO2 precursor
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Figure 3. HR-TEM images with SAED pattern (inset) of the hydrothermally lithiated (a) δ-MnO2 and (b) R-MnO2, with hexagonal spots (SAED) highlighted.
Figure 4. N2 adsorption-desorption isotherms of (a) hydrothermally lithiated δ-MnO2 and (b) hydrothermally lithiated R-MnO2. The close and open symbols represent adsorption and desorption data, respectively. The inset indicates the pore size distribution curves.
possesses 2 × 2 tunnel 1D structure, which is quite different from the 2D layered structure of the lithiated Li1+xMnO3-δ. Therefore, the phase transition between these structures requires severe rearrangements of manganese and oxygen ions, resulting in the breakdown of the 1D nanowire into isolated 0D nanocrystals. The observed crystal morphology of the hydrothermally lithiated samples was confirmed by HR-TEM analysis (see Figure 3). We have also examined the local atomic arrangement of the lithiated materials with selected area electron diffraction (SAED) analysis. The hydrothermally lithiated R-MnO2 phase shows the hexagonal SAED pattern, characteristic of layered structure, indicating the single crystalline nature of the obtained 0D nanoparticles. On the contrary, we observed a diffuse ringtype SAED pattern from the hydrothermally lithiated δ-MnO2, which is due to its hierarchical superstructure composed of many differently oriented nanoplates. ICP and N2 Adsorption-Desorption Isotherm Measurements. The chemical composition of the hydrothermally lithiated samples was determined with ICP emission spectroscopy. The ICP analysis showed that the lithiated δ-MnO2 and R-MnO2 compounds have a high Li content with Li/Mn ratios of 1.70 and 1.51, respectively. This result cross-confirms the formation of lithium-rich Li1+xMnO3-δ phases after the hydrothermal LiOH treatment. We have investigated the surface area and porosity of the lithiated Li1+xMnO3-δ samples by measuring N2 adsorptiondesorption isotherms. As plotted in Figure 4, the hierarchical assembly of Li1+xMnO3-δ nanoplates shows a distinct hysteresis in the p/p0 > 0.6 region, indicating the presence of mesopores possibly formed by the porous stacking of component nanoplates.31,32 Such a hysteresis is not observed for the 0D Li1+xMnO3-δ nanoparticles because of the absence of porous
Figure 5. Mn K-edge XANES spectra for the precursor δ-MnO2 (thin solid lines) and its hydrothermally lithiated derivative (thick solid lines), the precursor R-MnO2 (thin dashed lines) and its hydrothermally lithiated derivative (thick dashed lines), and the references of Mn2O3 (circles) and β-MnO2 (triangles).
superstructure. A negligible adsorption of nitrogen in the low p/p0 region indicates the lack of existence of micropores in these materials. According to the fitting analysis with the BET equation, both the lithiated materials have the expanded surface area of ∼113 m2/g for the hierarchical assembly of 2D Li1+xMnO3-δ nanoplates and ∼43 m2/g for the 0D Li1+xMnO3-δ nanoparticles. The expanded surface area of the hydrothermal products underscores the overall formation of nanostructured and/or porous materials. The pore size distribution curves of the hydrothermally lithiated δ-MnO2 and R-MnO2 were calculated by the BJH (Barrett, Joyner, and Halenda) method with desorption branch data (inset of Figure 4).33 The hydrothermally lithiated δ-MnO2 was found to possess mesopores with an average diameter of ∼4.6 nm, whereas there is no mesopore in the hydrothermally lithiated R-MnO2. Mn K-Edge XANES Analysis. The local crystal structure and electronic configuration of manganese ions in the precursor manganates and their hydrothermally lithiated derivatives have been investigated using X-ray absorption near-edge structure (XANES) analysis at the Mn K-edge. The Mn K-edge XANES spectra of the precursors δ-MnO2 and R-MnO2, and the hydrothermally lithiated Li1+xMnO3-δ are illustrated in Figure 5, together with the reference spectra of Mn2O3 and MnO2. The edge positions of both the lithiated derivatives are almost the same as that of the reference β-MnO2, indicating that these materials possess the tetravalent oxidation state of the manganese ion. This confirms the formation of the lithium-rich Li1+xMnO3-δ phase since no marked decrease in Mn valency occurs after the
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Figure 7. Fourier-filtered Mn K-edge EXAFS spectra for (a) the precursor δ-MnO2 and (b) its hydrothermally lithiated derivative, and (c) the precursor R-MnO2 and (d) its hydrothermally lithiated derivative. The circles and solid lines represent the experimental and fitted data, respectively.
Figure 6. (Top) k3-Weighted Mn K-edge EXAFS spectra and (bottom) FT data for (a) the precursor δ-MnO2 and (b) its hydrothermally lithiated derivative, and (c) the precursor R-MnO2 and (d) its hydrothermally lithiated derivative. In the bottom panel, the circles and solid lines represent the experimental and fitted data, respectively.
lithiation reaction. In the case that LiMn2O4 phase is formed by the lithiation of the MnO2 precursor, a remarkable lowering of the Mn oxidation state should occur. This is not the case. As illustrated in Figure 5, all of the materials under investigation display weak pre-edge peaks P and P′, corresponding to the 1s f 3d transitions.27,34 Since these transitions are not allowed for centrosymmetric octahedral symmetry, the observed weak intensity of these features indicates the stabilization of Mn ions in the octahedral sites.35 In main-edge region, there are peaks A, B, and C related to dipole-allowed 1s f 4p transitions. The intensity and sharpness of peak C are known to be proportional to the relative concentration of edge-sharing over corner-sharing of MnO6 octahedra.27,34 While the precursor R-MnO2 with corner- and edge-shared MnO6 octahedra exhibits a rather broad and diffuse peak C, a sharp and intense peak C is observed for the lithiated derivative of R-MnO2. This observation highlights the fact that the lithiation of the R-MnO2 precursor results in the stabilization of manganese ions in the layered structures composed of edgeshared MnO6 octahedra only. In the case of the precursor δ-MnO2, a sharp and intense peak C is commonly observable for the pristine and lithiated samples, confirming the maintenance of the layered structure upon the lithiation process. Mn K-Edge EXAFS Analysis. The effect of the lithiation reaction on the local atomic arrangement around manganese ions in the precursor manganates has been probed quantitatively with the EXAFS technique. The top panel of Figure 6 represents the k3-weighted Mn K-edge EXAFS spectra for the precursors δ-MnO2 and R-MnO2, and their hydrothermally lithiated derivatives. The overall EXAFS oscillation of the δ-MnO2 precursor remains nearly the same upon the hydrothermal LiOH treatment, suggesting the maintenance of a layered structure. In contrast, the hydrothermal LiOH treatment induces a distinct change in
the EXAFS data of the R-MnO2 precursor with the notable depression of oscillation amplitude, reflecting a significant modification of local structure around manganese ions. A closer inspection reveals that the hydrothermally lithiated R-MnO2 shows a merged single peak around k ) 7-8 like the pristine and hydrothermally lithiated δ-MnO2 samples, which is quite distinguishable from the split features of the precursor R-MnO2 in the same k region. Despite significant reduction of oscillation amplitude, the overall spectral feature of the lithiated R-MnO2 is similar to that of layered δ-MnO2, suggesting the structural transition from the R-MnO2 structure to the layered structure upon hydrothermal LiOH treatment. The corresponding Fouriertransformed (FTs) spectra are illustrated in the bottom panel of Figure 6. There are three peaks at ∼1.6, ∼2.6, and ∼3.2 Å in the FT data of precursor R-MnO2, which are assigned as the (Mn-O) shell, edge-shared (Mn-Mn) shell, and corner-shared (Mn-Mn) shell, respectively. In contrast, the precursor δ-MnO2 shows two FT peaks corresponding to the (Mn-O) and edgeshared (Mn-Mn) shells, but no FT peak of corner-shared (Mn-Mn) shell appears beyond ∼3 Å. This is consistent with its layered structure consisting of edge-shared MnO6 octahedra only.36 Similarly, only two FT features can be observed commonly for the hydrothermally lithiated derivatives. This can be regarded as clear evidence for the layered structure of the lithiated products composed of edge-shared MnO6 octahedra only. In comparison with the precursor δ-MnO2, the hydrothermally lithiated derivatives exhibit somewhat lower intensity for the FT peak of the edge-shared (Mn-Mn) shell, suggesting minute difference in their crystal structures. It has been well known that the intralayer octahedral sites of the layered δ-MnO2 lattice are fully occupied by Mn ions, and hence, the Mn ions in this structure have six neighboring Mn ions linked by edge-sharing of octahedra.28 In contrast, the intralayer octahedral sites of the Li-rich Li2MnO3 lattice are shared by Mn and Li ions, leading to the decrease of the coordination number (CN) for edge-shared (Mn-Mn) shells to 3. In this regard, the depression of the (Mn-Mn) shell after the hydrothermal lithiation reaction strongly suggests the formation of the lithium-rich lithium manganate phase having a lower CN for edge-shared (Mn-Mn) shells. In order to quantitatively determine the local structural parameters, the FT peaks were inversely Fourier transformed to k space and curve-fitted. As presented in Figure 7, the present Mn K-edge EXAFS spectra of the precursors are quite repro-
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TABLE 1: Results of Curve Fitting Analysis for the Mn K-Edge EXAFS Spectra sample
bond
CN
R (Å)
σ2 (10-3 × Å2)
precursor δ-MnO2a
(Mn-O) (Mn-Mn)edge (Mn-O) (Mn-Mn)edge (Mn-Mn)corner (Mn-O) (Mn-Mn)edge (Mn-O) (Mn-Mn)edge
6.0 6.0 6.0 4.0 4.0 6.0 3.0 6.0 3.0
1.90 2.84 1.90 2.87 3.44 1.89 2.84 1.90 2.79
4.46 6.97 3.84 4.47 5.21 6.18 5.12 4.83 8.11
precursor R-MnO2b lithiated δ-MnO2c lithiated R-MnO2d
a The curve fitting analysis was performed for R ) 1.012-2.853 Å and k ) 3.85-11.05 Å-1. b The curve fitting analysis was performed for R ) 0.951-3.375 Å and k ) 3.90-12.35 Å-1. c The curve fitting analysis was performed for R ) 1.012-2.853 Å and k ) 3.75-11.05 Å-1. d The curve fitting analysis was performed for R ) 0.951-2.792 Å and k ) 3.95-12.30 Å-1.
Figure 8. Discharge capacity plots of hierarchically assembled δ-MnO2 2D nanoplates (2), R-MnO2 1D nanowires (b), hierarchically assembled Li1+xMnO3-δ 2D nanoplates (4), Li1+xMnO3-δ 0D nanoparticles (O), and bulk Li2MnO3 (0).
ducible with the models of δ-MnO2- and R-MnO2-type structures, respectively.34 As summarized in Table 1, the fitting analyses gave reasonable structural parameters for each crystal structure. The manganese ions in the δ-MnO2-structured compounds are coordinated by six oxygen ligands at 1.90 Å and six manganese ions at 2.84 Å, whereas there are three types of neighbors around manganese ions in the R-MnO2-structured compounds, i.e., six oxygen ligands at 1.90 Å, four manganese neighbors at 2.87 Å, and four manganese neighbors at 3.44 Å. In the cases of the hydrothermally lithiated derivatives, both the EXAFS data could be well-fitted with layered Li2MnO3 structure having a CN of 3 for the (Mn-Mn) shell, confirming the formation of the lithium-rich lithium manganate phase after the hydrothermal LiOH treatment. The curve fitting analysis for the hydrothermally lithiated δ-MnO2 gave a reasonable amplitude reduction factor (S02) of ∼0.75, which is similar to the values obtained for both precursors. On the contrary, a marked decrease in this factor was observed for the hydrothermally treated R-MnO2, i.e., ∼50% of the S02 value of the precursors, reflecting the presence of significant local structural disorder in this material. Such a large disorder in crystal structure reflects the incomplete rearrangement of Mn and O ions during the lithiation process due to structural dissimilarity between the precursor R-MnO2 and its lithiated Li1+xMnO3-δ derivative. Electrochemical Measurement. The electrode performance of the lithiated Li1+xMnO3-δ materials was compared with those of the δ-MnO2 and R-MnO2 precursors, and bulk Li2MnO3, in order to probe the influences of the lithiation reaction and phase transformation on the electrochemical properties of precursor manganese oxides. The discharge capacities of the lithiated Li1+xMnO3-δ nanostructures are plotted as a function of cycle number in Figure 8, in comparison with those of the precursors δ-MnO2 and R-MnO2, and bulk Li2MnO3 microcrystals. Each lithium-rich Li1+xMnO3-δ material exhibits better electrode performance than the corresponding precursor manganese oxide and bulk Li2MnO3, highlighting the positive effects of phase transformation and nanostructure formation on the electrode performance of the nanostructured manganese oxides. The hierarchical assembly of 2D Li1+xMnO3-δ nanoplates begins with a high discharge capacity of 280 mAh/g, which drops to 223 mAh/g in the fifth cycle. Even with the following minute gradual decrease in capacity, this material delivers a large capacity of 165 mAh/g for the 50th cycle, which is much larger than the corresponding capacity of the precursor δ-MnO2 (118 mAh/g). The isolated 0D Li1+xMnO3-δ nanoparticle shows a large discharge capacity of 208 mAh/g for the first cycle and
140 mAh/g for the 50th cycle, which is much superior to the discharge capacity of the precursor R-MnO2 (169 mAh/g for the first cycle and 99 mAh/g for the 10th cycle). The observed capacity improvement upon lithiation is attributable to the fact that the phase transition from hydrated R-MnO2 or hydrated δ-MnO2 structure to nonhydrated layered structure facilitates the intercalation-deintercalation of lithium ions. Between the hydrothermally lithiated materials, the hierarchical assembly of Li1+xMnO3-δ nanoplates delivers larger discharge capacity for all of the present cycles because the component 2D nanoplates with a smaller thickness of ∼30 nm are more favorable for Li+ ion diffusion compared with 0D nanoparticles with a thicker diameter of ∼50 nm. In addition, as found from the Mn K-edge EXAFS analysis, the larger local structural disorder in the 0D nanoparticles is also responsible for their inferior electrode performance over the hierarchically assembled 2D nanoplates. Of note is that in comparison with the nanostructured Li-rich Li1+xMnO3-δ samples, the bulk homologue exhibits a much smaller discharge capacity of ∼30-40 mAh/g. This observation provides strong evidence for the advantage of nanostructure formation in improving the electrode performance of the Lirich lithium manganate phase. To probe the effects of lithiation on the redox behavior of the nanostructured manganese oxides, CV measurements were carried out for the precursor manganates and their hydrothermally lithiated derivatives. As plotted in Figure 9, the lithiation process leads to remarkable current increase for both the precursors δ-MnO2 and R-MnO2, which is consistent with the notable improvement of discharge capacity caused by lithiation (Figure 8). Between the lithiated materials, the nanoworm-like hierarchical sphere of Li1+xMnO3-δ delivers a larger current than the Li1+xMnO3-δ nanoparticle, which agrees well with the larger discharge capacity of the former. Overall features of CV data are nearly identical for both lithiated compounds since these materials have the same crystal structure and very similar chemical composition. As presented in Figure 9, a distinct oxidation peak appears at ∼3.4 V, whereas two reductions peaks are observed at ∼3.8 and ∼2.5 V. The oxidation and reduction peaks at higher potential of ∼3.4 and ∼ 3.8 V correspond to the deintercalation of lithium ions with the oxidation of Mn3+ to Mn4+ and to the lithium intercalation into the bulk sites with the reduction of Mn4+ to Mn3+, respectively. The other reduction peak at a lower potential of ∼2.5 V could be interpreted as the intercalation of lithium ions into surface sites since this process is known to occur at lower potential compared with the lithium intercalation into the bulk state.18 All of the present materials
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J. Phys. Chem. C, Vol. 113, No. 40, 2009 17397 On the contrary, the crystal structures of the lithiated nanostructured materials remain unchanged after electrochemical cycling. This finding provides straightforward evidence for the improvement of the structural stability of the precursor manganates by the lithiation reaction. Such an enhanced stability of the lithiated Li1+xMnO3-δ samples is responsible for the observed improvement of the electrode performance after the lithiation process. Conclusions
Figure 9. CV data of the precursors (a) δ-MnO2 and (b) R-MnO2, and the hydrothermally lithiated (c) δ-MnO2 and (d) R-MnO2 for the first and the second cycles.
Figure 10. Powder XRD patterns of manganese oxides after the 50th charge-discharge cycles: the precursors (a) δ-MnO2 and (c) R-MnO2, and (b,d) their hydrothermally lithiated derivatives. The asterisks and circles represent the superlattice ordering of the Li2MnO3 phase and the graphite conductor, respectively. In c, the underlined miller indices denote the Bragg reflections of the R-MnO2 phase.
exhibit no remarkable variations of CV data for the first cycle and the following cycles, strongly suggesting the good stability of the redox process in these materials. XRD Analysis for Electrochemically Cycled Derivatives. The structural variation of the present manganese oxides after the 50th cycle was examined with powder XRD analysis. As plotted in Figure 10, the precursor δ-MnO2 experiences a distinct phase transformation into layered Li2MnO3 or spinel-structured LiMn2O4 structure. Although a significant lowering of crystallinity caused by the cycling prevents us from conclusively determining the exact structure of the electrochemically cycled sample between layered or spinel structures, it is obvious that the electrochemical cycling causes severe structural frustration of the precursor δ-MnO2. In the case of the electrochemically cycled R-MnO2 precursor (Figure 10c), the underlined miller indices denote the Bragg reflections of the R-MnO2 phase, whereas the miller indices without underlining denote those of the Li-rich Li2MnO3. The Bragg reflections of layered lithium manganate phases can be observed together with those of the precursor R-MnO2 structure, strongly suggesting the occurrence of structural transformation during the cycling as in the case of the precursor δ-MnO2.
The hydrothermal LiOH treatment for nanostructured manganese oxide precursors makes it possible to synthesize the nanoworm-like hierarchical assembly of 2D nanoplates and the isolated 0D nanocrystals of the Li-rich Li1+xMnO3-δ phase. The resulting lithiated materials possess a Li-rich Li2MnO3type layered structure with the tetravalent Mn oxidation state and expanded surface area. The nanostructured Li1+xMnO3-δ materials exhibit better electrode performance compared with that of the precursor manganese oxides and bulk Li-rich lithium manganate. A phase transition to Li-rich Li1+xMnO3-δ with high structural stability is believed to be responsible for the improvement of the electrode performance of nanostructured manganese oxides upon the lithiation process. The present phase transformation method allows us to synthesize nanostructured lithiumrich manganese oxides with controllable dimensionality and morphologies, and to optimize the electrode performance of nanostructured manganese oxides as well. Currently, we are trying to apply the hydrothermal LiOH treatment for other types of nanostructured manganese oxides. Acknowledgment. This work was supported by Daegu Gyeongbuk Institute of Science & Technology funded by Ministry of Education, Science and Technology and partly by the Korea Science and Engineering Foundation (KOSEF) Grant funded by the Korea government (MEST) (grant: 20080061493) and the SRC/ERC program of MOST/KOSEF (grant: R11-2005-008-03002-0). The experiments at Pohang Accelerator Laboratory (PAL) were supported in part by MOST and POSTECH. References and Notes (1) Thackeray, M. M.; Rossouw, M. H.; Kock, A.; Harpe, A. P. J. Power Sources 1993, 43-44, 289. (2) Thackeray, M. M. Mater. Res. Bull. 1983, 18, 461. (3) Zhu, S. M.; Zhou, H. S.; Hibino, M.; Honma, I.; Ichihara, M. AdV. Funct. Mater. 2005, 15, 381. (4) Armstrong, A. R.; Bruce, P. G. Nature 1996, 381, 499. (5) Thackeray, M. M.; Johnson, P. J.; Depicciotto, L. A.; Bruce, P. G.; Goodenough, J. B. Mater. Res. Bull. 1984, 19, 179. (6) Thackeray, M. M.; Dekock, A. J. Solid State Chem. 1988, 74, 414. (7) Luo, J.-Y.; Wang, Y.-G.; Xiong, H.-M.; Xia, Y.-Y. Chem. Mater. 2007, 19, 4791. (8) Thackeray, M. M.; Johnson, C. S.; Vaughey, J. T.; Li, N.; Hackney, S. A. J. Mater. Chem. 2005, 15, 2257. (9) Lu, Z.; Dahn, J. R. J. Electrochem. Soc. 2002, 149, A1454. (10) Lu, Z.; Dahn, J. R. J. Electrochem. Soc. 2002, 149, A815. (11) Hwang, S.-J.; Park, H. S.; Choy, J.-H.; Campet, G. J. Phys. Chem. B 2000, 104, 7612. (12) Reed, J.; Ceder, G. Chem. ReV. 2004, 104, 4513. (13) Ammundsen, B.; Paulsen, J. AdV. Mater. 2001, 13, 943. (14) Hosono, E.; Kudo, T.; Honma, I.; Matsuda, H.; Zhou, H. Nano Lett. 2009, 9, 1045. (15) Kim, D. K.; Muralidharan, P.; Lee, H.-W.; Ruffo, R.; Yang, Y.; Chan, C. K.; Peng, H.; Huggins, R. A.; Cui, Y. Nano Lett. 2008, 8, 3948. (16) Shaju, K. M.; Bruce, P. G. Chem. Mater. 2008, 20, 5557. (17) Luo, J.-Y.; Xiong, H.-M.; Xia, Y.-Y. J. Phys. Chem. C 2008, 112, 12051. (18) Okubo, M.; Hosono, E.; Kim, J.; Enomoto, M.; Kojima, N.; Kudo, T.; Zhou, H.; Honma, I. J. Am. Chem. Soc. 2007, 129, 7444.
17398
J. Phys. Chem. C, Vol. 113, No. 40, 2009
(19) Futaba, D. N.; Hata, K.; Yamada, T.; Hiraoka, T.; Hayamizu, Y.; Kakudate, Y.; Tanaike, O.; Hatori, H.; Yumura, M.; Iijima, S. Nat. Mater. 2006, 5, 987. (20) Reddy, A. L. M.; Shaijumon, M. M.; Gowda, S. R.; Ajayan, P. M. Nano Lett. 2009, 9, 1002. (21) Lee, Y.; Kim, M. G.; Cho, J. Nano Lett. 2008, 8, 957. (22) Park, D. H.; Lim, S. T.; Hwang, S.-J.; Yoon, C.-S.; Sun, Y.-K.; Choy, J.-H. AdV. Mater. 2005, 17, 2834. (23) Park, D. H.; Ha, H.-W.; Lee, S. H.; Choy, J.-H.; Hwang, S.-J. J. Phys. Chem. C 2008, 112, 5160. (24) Park, D. H.; Lim, S. T.; Hwang, S.-J. Electrochim. Acta 2006, 52, 1462. (25) Park, D. H.; Lim, S. T.; Lee, S. H.; Hwang, S.-J.; Yoon, Y. S.; Lee, Y.-H.; Choy, J.-H. AdV. Funct. Mater. 2007, 17, 2949. (26) Kim, M. G.; Jo, M.; Hong, Y.-S.; Cho, J. Chem. Commun. 2009, 218. (27) Hwang, S.-J.; Kwon, C. W.; Portier, J.; Campet, G.; Park, H. S.; Choy, J.-H.; Huong, P. V.; Yoshimura, M.; Kakihana, M. J. Phys. Chem. B 2002, 106, 4053.
Baek et al. (28) Thackeray, M. M.; Kang, S.-H.; Johnson, C. S.; Vaughey, J. T.; Benedek, R.; Hackney, S. A. J. Mater. Chem. 2007, 17, 3112. (29) Pouillerie, C.; Croguennec, L.; Delmas, C. Solid State Ionics 2000, 132, 15. (30) Lu, Z.; Beaulieu, L. Y.; Donaberger, R. A.; Thomas, C. L.; Dahn, J. R. J. Electrochem. Soc. 2002, 149, A778. (31) Sing, K. S. W.; Everett, D. H.; Haul, R. A. W.; Moscou, L.; Pierotti, R. A.; Rouquerol, J.; Siemieniewska, T. Pure Appl. Chem. 1985, 57, 603. (32) Allen, T. Particle Size Measurement, 5th ed.; Chapman and Hall: London, 1997. (33) Yin, S.; Uchida, S.; Fujishiro, Y.; Aki, M.; Sato, T. J. Mater. Chem. 1999, 9, 1191. (34) Hwang, S.-J.; Choy, J.-H. J. Phys. Chem. B 2003, 107, 5791. (35) Treuil, N.; Labrugere, C.; Menetrier, M.; Portier, J.; Campet, G.; Deshayes, A.; Frison, J. C.; Hwang, S.-J.; Song, S. W.; Choy, J.-H. J. Phys. Chem. B 1999, 102, 2100. (36) Wells, A.-F. Structural Inorganic Chemistry; Clarendon Press: Oxford, 1984.
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