Origin of Improved Electrochemical Activity of β-MnO2 Nanorods

Oct 21, 2009 - 1D nanorods/nanowires of manganese oxides with different crystal structures and morphologies were prepared and characterized to underst...
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Origin of Improved Electrochemical Activity of β-MnO2 Nanorods: Effect of the Mn Valence in the Precursor on the Crystal Structure and Electrode Activity of Manganates In Young Kim,†,‡ Hyung-Wook Ha,†,‡ Tae Woo Kim,† Younkee Paik,§ Jin-Ho Choy,† and Seong-Ju Hwang*,† Center of Intelligent Nano-Bio Materials (CINBM), Department of Chemistry and Nano Sciences, Ewha Womans UniVersity, Seoul 120-750, and Korea Basic Science Institute, 67 Kyungdae Jungmoonro, Pookgu, Daegu 702-701, Korea ReceiVed: September 4, 2009; ReVised Manuscript ReceiVed: September 30, 2009

1D nanorods/nanowires of manganese oxides with different crystal structures and morphologies were prepared and characterized to understand the influence of the Mn valence in the solid-state precursor on the electrochemical activity of these nanomaterials and to elucidate the mechanism responsible for the excellent activity of β-MnO2 nanorods as well. According to powder X-ray diffraction analyses, treating manganese oxide precursors that have an oxidation state of e+3 with persulfate ions under hydrothermal conditions yields manganese oxides with the β-MnO2 structure. In contrast, the use of a LiMn2O4 precursor with a higher Mn valence leads to the formation of the R-MnO2-structured manganese oxide. Electron microscopic studies clearly show a 1D nanorod-type morphology for the β-MnO2 material, whereas a 1D nanowire-type morphology with a higher aspect ratio is observed for the R-MnO2 material. The diameter of the β-MnO2 nanorods decreases as the Mn valence in the precursors becomes smaller. According to electrochemical measurements, the formation of nanorods dramatically improves the electrode performance of the β-MnO2 phase. This compares with a relatively weak performance enhancement for the R- and δ-MnO2 phases upon the nanowire formation. The optimum electrode property results from the smaller β-MnO2 nanorods prepared with the MnO precursor. 7Li magic angle spinning nuclear magnetic resonance spectroscopy clearly demonstrates that Li+ ions in the lithiated β-MnO2 phase are adsorbed mainly on the sample surface. On the basis of this finding, we attribute the improved electrode performance of the β-MnO2 nanorods to their expanded surface area. Introduction During the past several decades, lithium rechargeable batteries have become the main power source for mobile electronics and telecommunication devices.1 The use of lithium rechargeable batteries has expanded to high-power applications, such as electric vehicles and power tools, which require efficient electrode materials with high energy density and excellent rate characteristics. Compared with microcrystalline electrode materials, nanocrystalline materials have many advantages, including a larger discharge capacity and higher structural stability during electrochemical cycling.2–8 Intense research has focused on the synthesis of nanocrystalline metal oxides and their electrode applications.9–14 Among their advantages, 1D nanowires/nanorods of manganese oxides boast a low price, low toxicity, facile synthesis, and rich abundance of Mn elements.15–17 Manganese oxides with 1D nanowire-type morphology show very promising performance as electrodes for lithium ion batteries, and their electrochemical functionality can be improved by the partial substitution of Mn with other metal ions.18–20 In the previous studies, 1D nanowires/nanorods of manganese oxide were synthesized via a soft-chemical redox reaction of crystalline manganese oxide precursors.21–23 From the viewpoint of crystal growth, the chemical nature of the solid* To whom correspondences should be addressed. Phone: +82-2-32774370. Fax: +82-2-3277-3419. E-mail: [email protected]. † Ewha Womans University. ‡ These authors contributed equally to this work. § Korea Basic Science Institute.

state precursor is expected to have a significant influence on the crystal morphology and physicochemical properties of the resulting manganate. There are, however, no studies of the effects of the Mn oxidation state and chemical composition of the precursor on the crystal structure, morphology, and electrochemical property of the manganese oxide 1D nanowires/ nanorods. An understanding of the relationship between crystal morphology and electrochemical activity would allow the optimization of the electrode performance of low-dimensional nanocrystalline manganese oxides. In line with this, it would be very intriguing to study the correlation between the structure type of manganates and the effect of the formation of a 1D nanowire or 1D nanorod on their electrochemical properties. Also information about the binding site of lithium ions in the 1D manganese oxide nanorods would give insight for the mechanism responsible for their electrode activity. In the present study, we synthesized 1D nanowires and nanorods of manganese oxides through hydrothermal persulfate treatments of various solid-state precursors with different Mn oxidation states and dissimilar chemical compositions. The crystal structure, crystal morphology, and chemical bonding nature of the obtained manganese oxides were systematically characterized with microscopic and spectroscopic tools. Using these nanomaterials, we studied the effect of the crystal morphology on the functionality of the manganese oxide as a lithium intercalation electrode. Also we examined a correlation between the crystal structure and electrode activity in the manganate nanowires or nanorods with R-, β-, and δ-MnO2-

10.1021/jp908556h CCC: $40.75  2009 American Chemical Society Published on Web 10/21/2009

Electrochemical Activity of β-MnO2 Nanorods

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type structures. To understand the remarkable improvement in electrode performance of the β-MnO2 phase upon the nanorod formation, the Li+ binding sites in the lithiated nanorods were investigated with 7Li magic angle spinning nuclear magnetic resonance, or MAS NMR, spectroscopy. Experimental Section Preparation. The 1D nanowires and nanorods of manganese oxides were prepared by persulfate treatments of the solid-state precursors of MnO, Mn2O3, MnO2, LiMn2O4, and Li2Mn2O4 under hydrothermal conditions. All the precursors except lithium manganates were purchased from Sigma-Aldrich Co. and used without further purification. The cubic spinel LiMn2O4 was prepared by solid-state reaction of the stoichiometric mixture of Li2CO3 and Mn2O3. The tetragonal spinel Li2Mn2O4 was obtained by n-BuLi treatment of the cubic spinel LiMn2O4 at room temperature for 72 h. For synthesis of the manganate nanowires and nanorods, the precursors were reacted with a 0.5 M aqueous solution of (NH4)2S2O8 at 140 °C for 36 h in a Teflon-lined hydrothermal vessel. After the reaction, the hydrothermal vessel was cooled to room temperature. The obtained powders were washed thoroughly with distilled water and dried in an oven. The chemically lithiated derivatives of 1D β-MnO2 nanorods and bulk β-MnO2 were prepared by reaction with excess 1.6 M n-BuLi hexane solution at room temperature for 72 h. After the lithiation reactions, the obtained powders were thoroughly washed with excess hexane and dried in a vacuum. Characterization. The crystal structures of the manganese oxides were characterized by powder X-ray diffraction, or XRD (λ ) 1.5418 Å, 298 K), measurements. We used inductively coupled plasma spectrometry, or ICP, to determine the chemical composition of these samples. To estimate the average oxidation state of Mn in the samples, we have carried out the iodometric redox titration. Typically, an accurately weighed sample (ca. 35 mg) of the manganese oxide was dissolved in a mixed solution of 6 M HCl (ca. 100 mL) and excess KI (ca. 3 g). Generated neutral iodine (I2) in the solution was titrated with standardized sodium thiosulfate (Na2S2O3) solution (ca. 0.012 M). The morphology of the manganese oxides was probed by field emission scanning electron microscopy, or FE-SEM (Jeol JSM-6700F with an energy-dispersive X-ray spectrometer). To determine the surface area of the manganese oxides, N2 adsorption-desorption isotherms were measured volumetrically at 77 K after degassing of the samples at 150 °C for 2 h under a vacuum. The crystal dimension and local crystal structure of the manganese oxides were examined by high-resolution transmission electron microscopy/selected area electron diffraction, or HR-TEM/SAED (Jeol JEM-2100F, 200 kV), measurement. The X-ray absorption spectroscopy, or XAS, data at the Mn K-edge were collected from the thin layer of powder samples deposited on transparent adhesive tape in a transmission mode at beamline 7C at the Pohang Accelerator Laboratory (PAL; Pohang, Korea) operated at 2.5 GeV and 180 mA. The measurements were carried out at room temperature with a Si(111) single-crystal monochromator. All the spectra were calibrated by measuring the spectrum of Mn metal foil. The data analysis for the experimental spectra was performed by the standard procedure reported previously. Curve fitting was accomplished using the UWXAFS 2.0 program, with Fi(k), φi(k), and λ(k) theoretically calculated by a curved wave ab initio extended X-ray absorption fine structure, or EXAFS, code FEFF 6.01. In the course of fitting analysis for the reference bulk β-MnO2, the amplitude reduction factor (S02) was set as a variable and the coordination numbers (CNs) were fixed to the

Figure 1. Powder XRD patterns of the hydrothermally treated derivatives of the precursors (a) MnO, (b) Mn2O3, (c) MnO2, (d) LiMn2O4, and (e) Li2Mn2O4. The asterisk represents the (110) reflection of the γ-MnO2 phase.

crystallographic values. To curve fit the other samples, S02 was fixed to the value from the bulk β-MnO2; the CNs were set as variables to exactly estimate the CN. Also, the bond distances (R), Debye-Waller factors (σ2), and energy shifts (∆E) were allowed to vary. The reliability of the fit was evaluated by the F factor, F ) [∑{k3(χ(k)calcd - χ(k)exptl)}2/(n - 1)]1/2, where χ(k)exptl, χ(k)calcd, and n represent the experimental EXAFS oscillation, the fitted oscillation, and the number of data points, respectively. 7Li MAS NMR experiments were performed on a Varian 200 FT NMR spectrometer at the Korea Basic Science Institute. An MAS probe equipped with 2.5 mm rotors was used. The sample rotors were spun at 18-25 kHz. Spectra were recorded at 77.7 MHz using a rotor-synchronized Hahn-echo pulse sequence. The 90° pulse length was 3.0 µs, and the pulse repetition delay was 2 s. The 7Li MAS NMR spectra were referenced to that of a 1.0 M LiCl solution at 0 ppm. The isotropic resonance of an MAS pattern was determined by acquiring NMR spectra, varying the sample spinning. Electrochemical Measurement. The electrochemical measurements were performed with a cell of Li/1 M LiPF6 in ethylene carbonate/diethyl carbonate (EC:DEC ) 50:50, v/v)/ active material, which was assembled in a drybox. A 20 mg portion of active electrode material was mixed with 12 mg of conductive binder (i.e., 8 mg of teflonized acetylene black and 4 mg of graphite), and a pure lithium foil was used as the negative electrode. After assembly, the cells were stored at room temperature for 12 h to ensure complete impregnation of the electrodes and separators with the electrolyte solutions. All the experiments were carried out in a galvanostatic mode with a Maccor multichannel galvanostat/potentiostat in the voltage range of 1.0-4.4 V at constant current density. Results and Discussion Powder XRD and ICP Analyses. Figure 1 represents the powder XRD patterns of hydrothermally treated manganese oxides. We found that the β-MnO2 phase was commonly produced by hydrothermal persulfate treatments of the binary precursors MnO and Mn2O3. This indicates an increase in the Mn oxidation state during the hydrothermal reaction. Persulfate treatment of ternary LiMn2O4 precursors gives rise to the formation of the R-MnO2 phase with a minor γ-MnO2 phase.

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Figure 2. FE-SEM images of the hydrothermally treated derivatives of the precursors (a) MnO, (b) Mn2O3, (c) MnO2, (d) LiMn2O4, and (e) Li2Mn2O4.

This ternary precursor is different from the other binary precursors in that it has lithium ions as well as a higher Mn oxidation state of +3.5. Thus, we speculated that the formation of the R-MnO2-type structure could be related to either the stabilization of 2 × 2 pores in this structure by incorporated lithium ions or the higher Mn oxidation state of the LiMn2O4 precursor. To identify a main factor that would determine the crystal structure of the products, we also carried out hydrothermal persulfate treatment of tetragonal spinel Li2Mn2O4 because this compound contains lithium ions and the lower Mn valence state of +3 as well. As shown in Figure 1, the hydrothermal treatment of Li2Mn2O4 leads to the production of β-MnO2-type material, as in the cases of the other binary MnOx precursors with a Mn valence state of e+3. This finding emphasizes that the crystal structure of hydrothermally treated manganese oxides is strongly dependent on the average oxidation state of the manganese ion in the precursors, not on the presence of alkali-metal ions. This conclusion is further supported by ICP spectrometry, which shows negligible incorporation of lithium ions in R-MnO2 nanowires prepared from the LiMn2O4 precursor. Elemental analysis also showed that only a negligible amount of ammonium cations exist in this R-MnO2 material (i.e., N/Mn ) 0.1%). We believed that the R-MnO2 structure of this material is stabilized by the incorporation of hydronium (H3O+) ions into its tunnel sites, instead of that of ammonium or lithium ions. As shown in Figure 1, the persulfate treatment of the precursor β-MnO2 caused no notable change in its XRD pattern. This reflects the fact that this precursor material has a tetravalent Mn oxidation state and the β-MnO2 structure. Summarizing the present experimental findings, we found that the use of manganese oxide precursors with the Mn oxidation state of e+3 results in the formation of β-MnO2-structured nanorods.24 The R-MnO2-type nanowires with minor γ-MnO2 phase are obtained with the manganate precursor having the Mn valence state of +3.5. Also, we have examined the Mn valences of the hydrothermally treated manganese oxides with iodometric titration. The oxidation states of Mn ions in the β-MnO2 phases prepared with the precursors MnO and Mn2O3 were estimated as +3.94 and +4.00, respectively. The R-MnO2 phase synthesized from the precursor LiMn2O4 possesses the Mn oxidation state of +3.51. This finding demonstrates that the β-MnO2 phases show a higher Mn oxidation state than the R-MnO2 phase. FE-SEM and HR-TEM Analyses. The crystal morphology of the hydrothermally treated manganese oxides was probed with

FE-SEM. As illustrated in Figure 2, 1D nanorods with a diameter of ∼50-270 nm and a length of several micrometers were obtained through the persulfate treatment of MnO, Mn2O3, and Li2Mn2O4. However, the use of LiMn2O4 precursor produced nanowires with a higher aspect ratio. On the contrary, the MnO2 precursor did not undergo notable changes in its morphology upon hydrothermal treatment. This observation shows that the oxidation of Mn ions in the precursor is a prerequisite for the formation of manganate 1D nanowires and nanorods. The formation of the intermediate layered phase is very important in producing 1D nanowires and nanorods,16 since the 1D nanomaterials can be formed through the rolling and folding of the intermediate layered crystallites. In fact, we successfully synthesized β-MnO2-structured nanowires via the hydrothermal treatment of layered manganese oxide.19 The aspect ratio of those β-MnO2 nanowires is much higher than that of the nanorods synthesized in the present study (Figure 2). We believe that the use of the microcrystalline-layered LiMnO2 precursor in the previous studies was responsible for the higher aspect ratio of those β-MnO2 nanowires.19 In the present study, the nonlayered manganate precursors were dissolved and recrystallized as intermediate layered crystallites to produce the 1D nanowires or nanorods. Such a multistep process is less effective in forming 1D nanowires compared with the rolling process of pre-existing layered crystallites. One consequence is a smaller aspect ratio. Conversely, the R-MnO2-type 1D nanowires prepared from LiMn2O4 precursor show a higher aspect ratio in spite of the nonlayered structure of the precursor adopted. We attributed this result to the inherent growth pattern of the R-MnO2 structure, preferring anisotropic growth into 1D morphology along the direction of 2 × 2 tunnels. A closer inspection of the FE-SEM images reveals that the average diameter of the nanorods is somewhat larger for the nanorods obtained from the trivalent manganate precursors Mn2O3 and Li2Mn2O4 (∼180-270 nm) than for those prepared from the lower valent manganese oxide, MnO (∼50-140 nm). It is well-known that the growth of manganese oxide nanowires or nanorods in a hydrothermal vessel precedes the dissolution oftheprecursorandthesubsequentformationof1Dnanomaterials.16,25 In this regard, the observed differences in crystal dimension of the nanorods is related to the dissimilar solubility of the precursors used; the trivalent manganese ions in the oxide lattice yield soluble Mn2+ and insoluble Mn4+ ions disproportionately.26 Thus, the dissolution of trivalent manganate precursors leads to the formation of insoluble Mn4+-containing crystallites that can act as seeds for the subsequent growth of the 1D nanowires

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Figure 3. HR-TEM/SAED data of the hydrothermally treated derivatives of the precursors (a) MnO, (b) Mn2O3, (c) MnO2, (d) LiMn2O4, and (e) Li2Mn2O4.

or nanorods of manganese oxide. The presence of seeds allows us to account for the larger crystal size of the nanorods prepared with the trivalent manganese oxide precursors. However, the negligible formation of insoluble tetravalent seeds from the lower valent manganese oxide precursors retards the crystal growth of the nanorods, which is reflected in the smaller diameter. The dependence of the crystal size on the Mn valence of the precursorisfurthersupportedbytheresultsofN2 adsorption-desorption isotherm measurements. According to the fitting analysis based on the Brunauer-Emmett-Teller, or BET, equation, a decrease of Mn valency in the precursors leads to an increase in surface area, 32 m2/g for the precursor Mn2+O, 8 m2/g for the precursor Mn3+2O3, and 6 m2/g for the precursor Li2Mn3+2O4.24 Since the lengths of the present nanorods are similar to one another, the surface area of the nanorods is inversely proportional to their diameter. In this regard, we found that the observed dependence of the surface area on the Mn oxidation state of the precursor provides strong support for the FE-SEM results (Figure 2). From the present experimental findings, it can be concluded that the lower valent manganate precursor is useful in preparing manganese oxide nanorods with a smaller diameter and larger surface area. The crystal dimension and shape of hydrothermally treated manganese oxides were also probed with HR-TEM; see Figure 3. As in the FE-SEM data, the nanorods can be observed in the HR-TEM images of the hydrothermally treated MnO, Mn2O3, and Li2Mn2O4, whereas the more anisotropic morphology of the 1D nanowire characterizes the hydrothermally treated LiMn2O4. The observed size and shape of the manganate are in good agreement with estimates from the FE-SEM results. The increase in the diameter of the β-MnO2 nanorods with an increase in Mn valency in the precursors is confirmed by the present HR-TEM results. Also, formation of 1D nanowires or 1D nanorods is not verified with the precursor MnO2 by HRTEM analysis (Figure 3c). Each material exhibits characteristic SAED patterns that match well with the corresponding crystal structure (i.e., R- or β-MnO2 structure). Mn K-Edge XANES Analysis. We investigated the effect of nanowire or nanorod formation on the local atomic arrangement and electronic configuration of manganese ions in the precursors using Mn K-edge X-ray absorption near-edge structure, or XANES, analysis. In Figure 4, the Mn K-edge XANES spectra of the hydrothermally treated derivatives of various precursors are compared with those of the corresponding precursors. Except for MnO2, all of the precursors show the blue shift of the absorption edge caused by the persulfate

Figure 4. Mn K-edge XANES spectra of the precursor and hydrothermally treated (a) MnO, (b) Mn2O3, (c) MnO2, (d) LiMn2O4, and (e) Li2Mn2O4. The solid lines represent the data of the pristine manganese oxides and the dashed line those of their hydrothermally treated samples. The right panel provides expanded views of pre-edge spectra in the energy region of 6537-6545 eV.

treatment, highlighting the oxidation of Mn ions during hydrothermal treatment. On the contrary, no notable shift of the edge position was detected for the β-MnO2 precursor, which is in good agreement with the observed absence of change in morphology and crystal structure. In the pre-edge region of 6535-6545 eV, all materials in the present study exhibit weak peaks P and/or P′, corresponding to the 1s f 3d transitions. These peaks can be enhanced by the distortion of Mn local symmetry from centrosymmetric octahedra to noncentrosymmetric tetrahedra.27 All the Mn ions in the precursors exist in octahedral symmetry. The high intensity of peak P for the MnO precursor is a result of simple overlap between this feature and a main-edge jump. In addition to the local symmetry of manganese ions, the density of unoccupied Mn 3d orbitals also affects significantly the intensity of these pre-edge peaks since the intensity of the XANES feature is proportional to the density of unoccupied final states.28 As shown in the right panel of Figure 4, the most intense pre-edge features appear for the β-MnO2 precursor of highest Mn valence with the highest hole density in 3d orbitals. Upon the hydrothermal treatment, all the precursors except LiMn2O4 show remarkable enhancement of the pre-edge peaks, just like the reference β-MnO2, confirming the formation of β-MnO2-type materials. In the case of the R-MnO2 nanowires made from the LiMn2O4 precursor, the intensity of the higher

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Figure 5. Potential profiles of the fifth cycle of the hydrothermally treated derivatives of (a) MnO, (b) Mn2O3, and (c) LiMn2O4 compared with those of (d) bulk β-MnO2. The applied current condition is 20 mA/g.

energy peak P′ is somewhat weaker than that of the reference MnO2. Considering the fact that the relative intensity of peak P′ over peak P is proportional to the average oxidation state of the manganese ion,27 we believe a weaker intensity of peak P′ can be regarded as evidence of the mixed Mn3+/Mn4+ oxidation state in the nanowires. Such a dependence of the Mn oxidation state on the precursor is in good agreement with the results of iodometric titration. In the main-edge region, there are some peaks, denoted A, B, and C, related to the dipole-allowed 1s f 4p transitions, whose shapes reflect the local environment of manganese ions.28 After the hydrothermal treatment, all the samples except LiMn2O4 showed significant changes, with these main-edge features resembling β-MnO2-like features. This underscores the modification of the local atomic arrangement to β-MnO2-type structure. Electrochemical Measurement. We examined the functionalities of the 1D nanowires and nanorods of manganese oxides as lithium intercalation electrodes to understand the effect of the crystal morphology on the electrochemical properties of these manganese oxides. Figure 5 illustrates the potential profiles of the hydrothermally treated derivatives of various precursors. In the initial cycle, the β-MnO2 nanorod prepared with the MnO precursor shows a larger discharge capacity of 373 (mA h)/g than the other materials (see the Supporting Information). Compared with the R-MnO2-structured nanowires prepared from the precursor LiMn2O4 (∼160 (mA h)/g for the fifth cycle), the β-MnO2 nanorods display a larger discharge capacity (∼240-290 (mA h)/g for the fifth cycle) with a wider 3 V plateau. In addition, the working potential of the R-MnO2-structured nanowires is quite lower than that of the β-MnO2 nanorods. In fact, it was reported that the 1D nanowire of R-MnO2 is useful as an anode material in a Li secondary battery, instead of a cathode material.20 As plotted in Figure 5, the bulk β-MnO2 microcrystals are electrochemically inactive. Of special interest is that the β-MnO2 nanorods exhibit high electrochemical activity, highlighting the effectiveness of nanorod formation in providing the inactive β-MnO2 phase with electrode activity. To understand the redox behavior of the obtained β-MnO2 nanorods, we have calculated their differential capacity plots, i.e. dQ/dV vs potential (not shown here). It was found that all the β-MnO2-structured materials commonly show a main reduction peak with the strongest intensity at ∼2.9 V in the course of the discharge process. This intense reduction peak corresponds to the lithium insertion with the reduction of Mn4+ to Mn3+.

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Figure 6. Discharge capacities of the hydrothermally treated derivatives of the precursors (a) MnO, (b) Mn2O3, and (c) LiMn2O4 compared with those of (d) bulk β-MnO2. The applied current condition is 20 mA/g.

The discharge capacities of the hydrothermally treated manganates are plotted in Figure 6 as a function of the cycle number. For the present cycle region, the nanorods synthesized from the MnO precursor deliver the largest capacity with the best cycle characteristics. As shown in Figure 6, its discharge capacity becomes quite stable after the third cycle. For the 30th cycle, this material shows the largest discharge capacity, 264 (mA h)/g, which is slightly lower than the theoretical discharge capacity of β-MnO2 (308 (mA h)/g). The relative order of the discharge capacity among the present electrode materials is maintained up to the 50th cycle. Also, we have found that, even when the current density is increased 10-fold (i.e., 20 f 200 mA/g), the β-MnO2 nanorods prepared with the MnO precursor still show acceptable cathode performance (∼107 (mA h)/g for the 50th cycle) (see the Supporting Information). It is believed the high rate performance of this nanorod material can be further improved by surface coating with highly conductive materials such as carbon and/or by reducing its diameter. These findings provide strong evidence for the fact that the precursor MnO is a very suitable precursor for preparing β-MnO2 nanorods with excellent electrode performance. The present β-MnO2 nanorods show much better electrode performance than the previously reported hierarchical superstructure of the β-MnO2 nanorods,9 suggesting the isolated morphology of the nanorods is more advantageous than the hierarchical assembly. On the other hand, it becomes apparent that the β-MnO2 nanorods deliver a much larger discharge capacity than the bulk β-MnO2 microcrystals, confirming the remarkable enhancement of the electrode performance of the β-MnO2 phase upon nanorod formation. With this finding, we took interest in variations of the electrochemical performance of other manganese oxide phases with different crystal structures upon nanorod formation. To address this issue, we prepared the 1D nanowires of manganese oxides with layered δ-MnO2 and tunnel R-MnO2 structures and their bulk counterparts as well (see the Supporting Information). We measured their performance as a lithium intercalation electrode. As can be seen from Figure 7, we found that, in contrast to the R-MnO2 phase showing only negligible variation of the discharge capacity after the nanowire was formed, the δ-MnO2 nanowire delivers a notably larger discharge capacity than its bulk counterpart. Even in the case of the δ-MnO2 phase, however, the effect of nanowire formation appears to be much less prominent compared with that of the β-MnO2 phase. The fact

Electrochemical Activity of β-MnO2 Nanorods

Figure 7. Discharge capacities of (a) R-MnO2 nanowires (squares), (b) bulk R-MnO2 microcrystals (triangles), (c) δ-MnO2 nanowires (circles), and (d) bulk δ-MnO2 microcrystals (tilted squares).

Figure 8. Powder XRD patterns of (a) β-MnO2 nanorods, (b) chemically lithiated β-MnO2 nanorods, (c) bulk β-MnO2 microcrystals, and (d) chemically lithiated bulk β-MnO2 microcrystals. In (b) and (d), Miller indices of Bragg reflections corresponding to the tetragonal spinel phase are underlined.

that only the β-MnO2 phase experiences dramatic improvement in the electrode property upon 1D nanorod formation raises special interest concerning the mechanism responsible for the electrochemical activity of this material. Effect of Lithiation on the Crystal Structure of β-MnO2 Nanorods. To elucidate the mechanism responsible for the excellent electrochemical activity of the β-MnO2 nanorods, we investigated the binding sites for the Li+ ions in the lithiated β-MnO2 nanorods and accompanying structural evolution with the XRD, XANES, and 7Li MAS NMR tools. The powder XRD patterns of lithiated β-MnO2 nanorods and microcrystals are plotted in Figure 8 with those of the corresponding pristine compounds. After the n-BuLi treatment, most of the Bragg reflections of the β-MnO2 nanorods become very diffuse and depressed, indicating amorphization of the β-MnO2 structure. Compared with the β-MnO2 nanorod, the bulk β-MnO2 material experiences relatively weak variations in the powder XRD pattern upon lithiation. This indicates a lower frustration of the crystal structure. To probe the local structural evolution of the β-MnO2 nanorods upon lithiation, we have carried out Mn K-edge XAS analysis. In Figure 9, the Mn K-edge XANES spectra of the pristine and lithiated β-MnO2 nanorods and β-MnO2 microc-

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Figure 9. Mn K-edge XANES spectra of the pristine and chemically lithiated (a) β-MnO2 nanorods, (b) bulk β-MnO2 microcrystals, and (c) spinel LiMn2O4. The solid lines represent the data of the pristine manganese oxides and the dashed lines those of their lithiated derivatives. The right panel provides expanded views of pre-edge spectra in the energy region of 6538-6546 eV.

rystals are compared with those of the cubic spinel LiMn2O4 and lithiated spinel Li2Mn2O4. For all the present manganese oxides, the n-BuLi treatment induces a distinct displacement of edge jump toward the low-energy side, underscoring the reduction of Mn ions. In the pre-edge region, there are weak features P and/or P′ related to the dipole-forbidden 1s f 3d transitions. The lithiated β-MnO2 nanorods and lithiated tetragonal spinel Li2Mn2O4 exhibit a suppression of peak P′ with a red shift of peak P; however, a distinct peak P′ is observable for the lithiated β-MnO2 microcrystals, as can be seen in the right panel of Figure 9. Since the relative intensity of peak P′ over peak P sensitively reflects the oxidation state of manganese ions, the incomplete suppression of the pre-edge peak P′ upon lithiation of the β-MnO2 microcrystals indicates the imperfect reduction of tetravalent Mn ions. Considering that the identical condition of n-BuLi treatment is applied to the β-MnO2 nanorods and microcrystals, we concluded that the present finding confirms that the nanorods are more susceptible to the Mn reduction caused by lithiation, which is related to their larger surface area. We investigated quantitatively the effect of lithiation on the local atomic arrangement of the β-MnO2 phase with EXAFS analysis. The Fourier-transformed, or FT, Mn K-edge EXAFS spectra of the pristine and lithiated β-MnO2 nanorods and microcrystals are plotted in the left panel of Figure 10 with those of the reference cubic spinel LiMn2O4 and tetragonal spinel Li2Mn2O4. The FT spectrum of the cubic spinel LiMn2O4 appears to be severely modified by lithiation. A distinct splitting of the Mn-Mn peak at ∼2.6 Å upon lithiation indicates a symmetry change from cubic to tetragonal caused by the Jahn-Teller distortion of trivalent manganese ions. On the other hand, both the β-MnO2 microcrystals and nanorods commonly display the typical FT spectra of the β-MnO2 phase: The first peak at ∼1.5 Å corresponds to the Mn-O bonding pair, whereas the following peaks at ∼2.5 and ∼3.1 Å originate from Mn-Mn shells in the edge- and corner-shared MnO6 octahedra pairs. Despite the marked depression of Mn-Mn peaks, the lithiated bulk β-MnO2 microcrystal exhibits FT data that are similar to those of the pristine compound. However, the β-MnO2 nanorods demonstrate a more prominent depression of FT peaks with remarkable displacement of their positions upon lithiation,

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Figure 10. (Left) FT and (right) inverse FT Mn K-edge EXAFS spectra of (a) the bulk β-MnO2 microcrystal, (b) the lithiated β-MnO2 microcrystal, (c) the β-MnO2 nanorod, (d) the lithiated β-MnO2 nanorod, (e) cubic spinel LiMn2O4, and (f) tetragonal spinel Li2Mn2O4. The circles represent the experimental data and the solid lines the calculated data.

TABLE 1: Results of Nonlinear Least-Squares Curve Fittings for the Mn K-Edge EXAFS Spectra of β-MnO2 Microcrystals, β-MnO2 Nanorods, and Spinel LiMn2O4 and Their Lithiated Derivatives sample a

bulk β-MnO2

lithiated bulk β-MnO2b

nano-β-MnO2c

lithiated nano-β-MnO2d cubic spinel LiMn2O4e tetragonal spinel Li2Mn2O4f

bond

CN

R (Å)

σ2 (10-3 Å2)

Mn-O Mn-O Mn-Mn Mn-Mn Mn-O Mn-O Mn-Mn Mn-Mn Mn-O Mn-O Mn-Mn Mn-Mn Mn-O Mn-O Mn-O Mn-Mn Mn-O Mn-O Mn-Mn Mn-Mn

4.0 2.0 2.0 8.0 2.2 1.1 0.8 3.0 4.0 2.0 2.0 8.0 1.9 0.9 4.8 4.8 4.0 2.0 2.0 4.0

1.89 1.91 2.87 3.42 1.92 1.94 2.80 3.38 1.89 1.90 2.86 3.42 1.93 1.94 1.90 2.91 1.89 2.27 2.80 3.01

2.48 2.48 2.14 3.85 1.83 1.83 0.57 6.32 3.86 3.86 3.86 4.63 2.18 2.18 4.63 4.14 3.63 2.15 2.93 7.01

a The curve-fitting analysis was performed for the range of R ) 0.92-3.405 Å and k ) 4.05-11.80 Å-1. b The curve-fitting analysis was performed for the range of R ) 0.92-3.405 Å and k ) 4.05-11.80 Å-1. c The curve-fitting analysis was performed for the range of R ) 0.951-3.528 Å and k ) 3.90-11.10 Å-1. d The curve-fitting analysis was performed for the range of R ) 0.92-1.994 Å and k ) 3.85-11.25 Å-1. e The curve-fitting analysis was performed for the range of R ) 0.982-3.037 Å and k ) 3.80-10.60 Å-1. f The curve-fitting analysis was performed for the range of R ) 0.982-3.007 Å and k ) 3.70-12.35 Å-1.

strongly suggesting the occurrence of severe structural modification. To determine structural parameters such as the coordination number, CN, bond distance, R, and Debye-Waller factor, σ2, these FT spectra were isolated by inverse FT to a k space and then curve-fitted.29 The experimental and fitted FT and inverse FT spectra of β-MnO2 microcrystals and nanorods are represented with their lithiated derivatives in Figure 10. They are compared with those of the cubic spinel LiMn2O4 and tetragonal spinel Li2Mn2O4. On the basis of the cubic and tetragonal spinel structures, we reproduced the EXAFS spectra of the reference LiMn2O4 and Li2Mn2O4. As listed in Table 1, the best fit bond distances of the Mn-O and Mn-Mn shells are very consistent with the crystallographic values of these phases, confirming the reliability of the present fitting analysis. In the cases of the pristine β-MnO2 microcrystals and nanorods, good fits were

Figure 11. 7Li MAS NMR spectra of (a) chemically lithiated β-MnO2 nanorods, (b) electrochemically lithiated β-MnO2 nanorods, (c) layered LiMn0.9Cr0.1O2, and (d) spinel LiMn2O4

commonly obtained with the β-MnO2 structure. Similarly, the EXAFS data of the lithiated bulk β-MnO2 microcrystals were well-fitted with the same β-MnO2 structure, clearly demonstrating the maintenance of the β-MnO2 structure upon lithiation. As listed in Table 1, the Mn-O bond distances of β-MnO2 microcrystals are slightly elongated by the lithiation process. This indicates the reduction of manganese ions. Also, the CNs of the lithiated β-MnO2 microcrystals are smaller than those of the pristine compound, strongly suggesting frustration of the atomic ordering around manganese ions. The β-MnO2 structure does not allow us to reproduce the experimental data of the lithiated β-MnO2 nanorods, highlighting the severe structural change of this type of nanorod upon lithiation. Instead we were able to reproduce the FT spectra of the lithiated β-MnO2 nanorods in the 0.9-2.0 Å region with the model of two Mn-O shells. The present fitting results provide straightforward evidence of the fact that the degree of structural modification upon lithiation was enhanced by nanorod formation. 7 Li MAS NMR Analysis. The binding sites of the lithium ions in the lithiated β-MnO2 phase were investigated with 7Li MAS NMR for electrochemically and chemically lithiated β-MnO2 nanorods. The electrochemically lithiated nanorods were separated from the composite cathode after the third electrochemical cycling. As plotted in Figure 11, the reference spinel LiMn2O4 exhibits a major resonance at 495 ppm, which is attributed to a lithium ion in the normal tetrahedral 8a site of the spinel structure.30 In addition, two additional resonances are clearly observed at 529 and 568 ppm, which correspond to the interstitial lithium ions in the octahedral 16c and 16d sites. Of special interest is that both the electrochemically and chemically lithiated β-MnO2 display the main resonance at -2 ppm, with additional weak signals at 108 and 503 ppm. In fact, the resonance at -2 ppm was observed previously as the dominant 6 Li NMR signal from the ionic lithium on the surface at the initial stage (e40% discharges) of the reduction of β-phase bulk manganese dioxide in Li/MnO2 primary (nonrechargeable) batteries.31 Since the lithium ions stabilized in the manganate lattice have strong covalent bonds with lattice oxygen, these lithium ions show a marked NMR shift of several hundred parts per million (ppm) due to notable electron transfer between

Electrochemical Activity of β-MnO2 Nanorods lithium and oxygen ions. Conversely, the lithium ions on the surface lack covalent bonding with lattice oxygen ions and hence show more ionic character compared with the lithium ions incorporated into the metal oxide lattice. Hence, the surfaceadsorbed lithium ions display only a negligible NMR shift. In this regard, the observed NMR spectra for the lithiated β-MnO2 nanorods strongly suggest the stabilization of lithium ions on the surface of the β-MnO2 nanorods/microcrystals not in the manganate lattice. Additional weak signals are accounted for by a small amount of lithium ions incorporated into the β-MnO2 lattice, at 503 ppm, and into the tetragonal-phase Li2Mn2O4, at 108 ppm. The origin of the weak signal at 186 ppm observed in the chemically lithiated β-MnO2 is not clear. A small amount of lithium manganate impurity might be formed during the harsh reaction conditions with 1.6 M n-BuLi. Unlike the lithiated β-MnO2 phase, the layered LiMn0.9Cr0.1O2 shows significant NMR shifts, resulting in resonance peaks at 142 and 198 ppm.32 Also, the lithiated R-MnO2 phase with 2 × 2 channels can show NMR resonance at 155-185 ppm.33 This result provides strong evidence for the stabilization of inserted lithium ions in the manganate lattice of R-MnO2, or layered manganate. The present experimental results demonstrate that the lithium ions in the lithiated β-MnO2 phase are adsorbed mainly on the sample surface, which is not the case with the R-MnO2 and layered δ-MnO2 phases. This finding is evidenced by the fact that the 1 × 1 channel in the β-MnO2 phase is too narrow for lithium ions to migrate into the manganate lattice,31 which is responsible for the poor electrode performance of bulk β-MnO2. Taking this into consideration, we attribute the remarkably increased discharge capacity of the β-MnO2 nanorods to their greatly expanded surface area with the many surface sites for Li+ grafting. Also, the number of surface sites for Li+ adsorption is expected to be greater with the decrease in the diameter of the nanorods and the increased surface area. Accordingly, the smallest β-MnO2 nanorods prepared from the precursor MnO show the largest discharge capacity among the present nanorods (Figure 6). In addition, the surface adsorption of the lithium ions allows us to explain the fact that the lithiation process affects more significantly the crystal structure of the β-MnO2 nanorods with its larger surface area than it does the structure of the β-MnO2 microcrystals (Figure 8). Conclusion We have investigated the effects of the Mn valence in the solid-state precursor on the crystal structure, morphology, and electrode activity of the manganate 1D nanowires and nanorods. Also, the origin of the improved electrode activity of the β-MnO2 phase upon the nanorod formation was investigated. The powder XRD and FE-SEM/HR-TEM analyses clearly demonstrate that the crystal structure and morphology of the resulting 1D nanocrystalline manganese oxide depend strongly on the Mn oxidation state of the precursor manganates. The low-valent precursor MnO is found to be the most effective precursor for the production of thin β-MnO2 nanorods with excellent electrode performance. One-pot hydrothermal treatment of the commercially available MnO is a simple and valuable synthetic route to a highly promising intercalation electrode of 1D manganate nanorods. On the other hand, the improvement in electrode performance upon the formation of nanowires or nanorods is found to be much more prominent for electrochemically inactive β-MnO2 phases than for active R-MnO2 and δ-MnO2 phases. According to the 7Li MAS NMR and XAS studies, the Li+ ions in the lithiated β-MnO2 nanorods are mainly adsorbed on their surface. Thus, the increase in

J. Phys. Chem. C, Vol. 113, No. 51, 2009 21281 surface area upon the formation of 1D nanorods is mainly responsible for the improvement of the electrode performance of the electrochemically inactive β-MnO2 phase with no crystal sites available for Li+ intercalation. Our group is investigating the synthesis of the 1D nanorods of other metal oxides with electrochemical inactivity and their application as electrodes for lithium secondary batteries. Acknowledgment. This work was supported by a Korea Science and Engineering Foundation (KOSEF) grant funded by the Korean Government (MEST) (2008-0061493), by the General R/D Program of the Daegu Gyeongbuk Institute of Science and Technology (DGIST), by the MEST of the Republic of Korea, by National Research Foundation of Korea Grant funded by the Korean Government (20090063005), and by the National R&D programs of the Ministry of Science and Technology (MOST), Republic of Korea. The experiments at PAL were supported in part by MOST and POSTECH. Supporting Information Available: Detailed preparation condition of the 1D nanowires of R-MnO2 and δ-MnO2 materials and their microcrystalline homologues, potential profiles of the initial cycle for the hydrothermally treated manganese oxides, plot of the surface area of the nanorods vs Mn valence in the precursors, FE-SEM images of the reference manganese oxide nanowires with R-MnO2 and δ-MnO2 structure, and discharge capacity plot of the β-MnO2 nanorods prepared from MnO under high current density. These materials are available free of charge via the Internet at http://pubs.acs.org. References and Notes (1) Kang, K.; Meng, Y. S.; Bre´ger, J.; Grey, C. P.; Geder, G. Science 2006, 311, 977–980. (2) Poizot, P.; Laruelle, S.; Grugeon, S.; Dupont, L.; Tarascon, J.-M. Nature 2000, 407, 496–499. (3) Larcher, D.; Masquelier, C.; Bonnin, D.; Chabre, Y.; Masson, V.; Leriche, J.-B.; Tarascon, J.-M. J. Electrochem. Soc. 2003, 150, A133A139. (4) Armstrong, A. R.; Canales, G.; Canales, G. J. R.; Bruce, P. G. AdV. Mater. 2005, 17, 862–865. (5) Arico, A. S.; Bruce, P.; Tarascon, J.-M.; Schalkwijk, W. V. Nat. Mater. 2005, 4, 366–377. (6) Singhal, A.; Skandan, G.; Amatucci, G.; Badway, F.; Ye, N.; Manthiram, A.; Ye, H.; Xu, J. J. J. Power Sources 2004, 129, 38–44. (7) Dobley, A.; Ngala, K.; Yang, S. F.; Zavalij, P. Y.; Whittingham, M. S. Chem. Mater. 2001, 13, 4382–4386. (8) Li, N.; Mitchell, C. D.; Lee, T. K. P.; Martin, C. R. J. Electrochem. Soc. 2003, 150, A979-A984. (9) Cheng, F.; Zhao, J.; Song, W.; Li, C.; Ma, H.; Chen, J.; Shen, P. Inorg. Chem. 2006, 45, 2038–2044. (10) Li, J.; Tang, Z.; Zhang, Z. Chem. Mater. 2005, 17, 5848–5855. (11) Xia, Y.; Yang, P.; Sun, Y.; Wu, Y.; Mayers, B.; Gates, B.; Yin, Y.; Kim, F.; Yan, H. AdV. Mater. 2003, 15, 353–389. (12) Chen, X.; Li, X.; Jiang, Y.; Shi, C.; Li, X. Solid State Commun. 2005, 136, 94–96. (13) Wang, X.; Li, Y. J. Am. Chem. Soc. 2002, 124, 2880–2881. (14) Langley, C. E.; Sljukic, B.; Banks, C. E.; Compton, R. G. Anal. Sci. 2007, 23, 165–170. (15) Yuan, J.; Li, W.-N.; Gomez, S.; Suib, S. L. J. Am. Chem. Soc. 2005, 127, 14184–14185. (16) Wang, X.; Li, Y. Chem.sEur. J. 2003, 9, 300–306. (17) Batchelor-McAuley, C.; Shao, L.; Wildgoose, G. G.; Green, M. L. H.; Compton, R. G. New J. Chem. 2008, 32, 1195–1203. (18) Hwang, S.-J.; Park, H. S.; Choy, J.-H.; Campet, G. J. Phys. Chem. B 2000, 104, 7612–7618. (19) Park, D. H.; Ha, H.-W.; Lee, S. H.; Choy, J.-H.; Hwang, S.-J. J. Phys. Chem. C 2008, 112, 5160–5164. (20) Li, B.; Rong, G.; Xie, L.; Huang, L.; Feng, C. Inorg. Chem. 2006, 45, 6404–6410. (21) 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, 2949–2956. (22) Park, D. H.; Lim, S. T.; Hwang, S.-J. Electrochim. Acta 2006, 52, 1462–1466.

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(23) Park, D. H.; Lim, S. T.; Hwang, S.-J.; Yoon, C. -S.; Sun, Y. K.; Choy, J.-H. AdV. Mater. 2005, 17, 2834–2837. (24) It was found that the hydrothermal persulfate treatment for Mn3O4 (the Mn valency is +2.67) also produced β-MnO2 nanorods with a surface area of 23 m2/g. This finding provided further support for our conclusions about the relationship between Mn valency in the precursor and the crystal structure/surface area of the hydrothermally obtained nanowires or nanorods. (25) Subramanian, V.; Zhu, H.; Vajtai, R.; Ajayan, P. M.; Wei, B. J. Phys. Chem. B 2005, 109, 20207–20214. (26) Deng, B.; Nakamura, H.; Yoshio, M. Electrochem. Solid-State Lett. 2005, 8, A171-A174. (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, 4053–4060. (28) Hwang, S.-J.; Choy, J.-H. J. Phys. Chem. B 2003, 107, 5791–5796.

Kim et al. (29) The best fitted residual F2 factor () ∑{k3(χ(k)calcd-χ(k)exptl)}2/(n1), where n is the number of data) was determined to be 0.052 for the bulk β-MnO2 microcrystals, 0.103 for the lithiated bulk β-MnO2 microcrystals, 0.010 for the β-MnO2 nanorods, 0.044 for the lithiated β-MnO2 nanorods, 0.074 for LiMn2O4, and 0.011 for Li2Mn2O4. (30) Grey, C. P.; Dupre, N. Chem. ReV. 2004, 104, 4493–4512. (31) Bowden, W.; Grey, C. P.; Hackney, S.; Wang, F.; Paik, Y.; Iltchev, N.; Sirotina, R. J. Power Sources 2006, 153, 265. (32) Pan, C.; Lee, Y. J.; Ammundsen, B.; Grey, C. P. Chem. Mater. 2002, 14, 2289–2299. (33) Tsuji, M.; Paik, Y.; Grey, C. P.; Murao, S. Solid-State Chem. Inorg. Mater. 2003, 407–412.

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