New Insight into Structural Evolution in Layered NaCrO2 during

Dec 8, 2014 - ... O3- to P′3-type phase through O′3-type phase by gliding of CrO2 ...... Shou-Hang Bo , Xin Li , Alexandra J. Toumar , and Gerbran...
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New Insight into Structural Evolution in Layered NaCrO2 during Electrochemical Sodium Extraction Kei Kubota,†,‡ Issei Ikeuchi,† Tetsuri Nakayama,† Chikara Takei,† Naoaki Yabuuchi,†,‡ Hiromasa Shiiba,§ Masanobu Nakayama,‡,§ and Shinichi Komaba*,†,‡ †

Department of Applied Chemistry, Tokyo University of Science, 1-3 Kagurazaka, Shinjuku, Tokyo 162-8601, Japan Elements Strategy Initiative for Catalysts and Batteries, Kyoto University, 1-30 Goryo-Ohara, Nishikyo-ku, Kyoto 615-8245, Japan § Department of Materials Science and Engineering, Nagoya Institute of Technology, Gokiso, Showa-ku, Nagoya, Aichi 466-8555, Japan ‡

S Supporting Information *

ABSTRACT: Electrochemical properties and structural changes during charge for NaCrO2, whose structure is classified as α-NaFeO2 type layered polymorph (also O3type following the Delmas’ notation), are examined as a positive electrode material for nonaqueous Na-ion batteries. NaCrO2 delivers initial discharge capacity of 110 mAh g−1 at 1/20C rate in the voltage range of 2.5−3.6 V based on reversible Cr3+/Cr4+ redox without oxidation to hexavalent chromium ions, while the initial discharge capacity is only 9 mAh g−1 when cutoff voltage is set to 4.5 V. Results from exsitu X-ray diffraction, X-ray absorption spectroscopy, and DFT calculations reveal that the irreversible phase transition occurs after sodium extraction by charging over a voltage plateau at 3.8 V associated with the lattice shrinkage along the c-axis in the case of x > 0.5 in Na1−xCrO2, which originates from the migration of chromium ions from octahedral sites in CrO2 slabs to both tetrahedral and octahedral sites in interslab layer. The irreversible structural change would disturb sodium insertion into the damaged layer structure during discharge, resulting in the loss of reversibility as electrode materials. Reversible cycle range with stable capacity retention is, therefore, limited to the compositional range of 0.0 ≤ x ≤ 0.5 in Na1−xCrO2.

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INTRODUCTION Rechargeable Na-ion batteries, which consist of Na-ion insertion materials, are promising candidates for the largescale energy storage applications because of the material abundance and availability of sodium.1−6 After 2010, various intercalation materials for the Na-ion batteries have been extensively studied to realize a rechargeable Na-ion battery system.7−17 Electrode performance of layered sodium transition metal oxides, which was originally studied as positive electrode materials for Na batteries in 1980s,18−23 have been significantly improved by the latest battery technology developed in the Liion battery system during the past three decades.24−27 As one of examples for the improvement, our group has succeeded in demonstrating practical sodium batteries (Na-ion batteries) with the long-life cycle performance, which consist of sodium intercalation guests, layered sodium transition metal oxide, and non-graphitizable carbon.28 Among layered sodium transition metal oxides, NaCrO2, which has classified as α-NaFeO2-type layered structure, attracts attention as a potential positive electrode material for sodium batteries due to its excellent cycle performance and thermal stability.24,29 Recently, Nohira and Chen et al. also demonstrated excellent capacity retention of Na//NaCrO2 cells using © XXXX American Chemical Society

molten salt as electrolyte at an elevated temperature (90 °C).30,31 The crystal structure of NaCrO2 can be also described as O3type layered structure using Delmas’ notation.32 The letter “O” indicates a coordination environment (i.e., octahedral sites) of alkali metals and “3” represents the number of MeO2 (Me = transition metal) slabs in a hexagonal unit cell. The edge-shared NaO6 and MeO6 octahedra stack into alternate layers along [001] direction in the hexagonal unit cell with the cubic close packed (CCP) oxygen array. The electrochemical reversibility of Na extraction/insertion from/into NaCrO2 was first reported by Braconnier et al. in 1982.20 They demonstrated the electrochemical activity as a positive electrode material in Na cells with an aprotic electrolyte and reported the phase transition from the O3-type to P3-type phase by electrochemical sodium extraction. We and some research groups carefully have examined its phase transition mechanisms, and the study on structural changes during sodium extraction revealed that the phase transition occurs from O3-type (x = 0 Received: October 21, 2014 Revised: December 5, 2014

A

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The Journal of Physical Chemistry C in Na1−xCrO2) to P′3-type (x = 0.5) via O′3-type (0.1 ≤ x ≤ 0.3) phase on sodium extraction, in which prime symbol in O′3 and P′3 indicates a distorted monoclinic phase.31,33,34 Despite their structural similarity, LiCrO2 was known as electrochemically inactive in a Li cell.35,36 On the other hand, NaCrO2 in a Na cell delivers a reversible capacity of ca. 120 mAh g−1 with good capacity retention. As a safety aspect for Na batteries, the reactivity between deintercalated Na1−xCrO2 and carbonate ester solution was examined using accelerating rate calorimetry (ARC), and the results revealed that Na0.5CrO2 has higher thermal stability compared with Li0.5CoO2 and Li0FePO4.29,37 Although NaCrO2 has a potential application as a positive electrode material for long-life and safe Na-ion batteries, reversible capacity of the O3-type NaCrO2 electrodes is relatively smaller than those of other O3-type layered positive electrode materials such as Na[Fe1/2Co1/2]O238 and Na[Fe0.4Ni0.3Mn0.3]O2.39 The reversible capacity for NaCrO2 is limited to ca. 120 mAh g−1, corresponding to ca. 0.5 mol of sodium extraction in Na1−xCrO2 and the reversibility is lost by further charge up to 4.5 V.33 Such loss of reversibility as the electrode material is presumably caused by an irreversible structural change when more than 0.5 mol of Na is extracted from the layered structure. To the best of our knowledge, no reports are found for the structural change during charge up to such a high voltage, and the detailed mechanism on the loss of reversibility for O3-type NaCrO2 is still unclear. In this study, to clarify the origin of capacity loss in relation to the structural changes, the crystal structures of Na1−xCrO2 (0.0 ≤ x ≤ 1.0) during charge up to 4.5 V have been examined by ex-situ X-ray diffraction measurements. Changes in electronic state of chromium atoms and their local structure during charge and discharge processes have also been studied by X-ray absorption spectroscopy. Furthermore, formation energies of O3- and P3type phases are estimated by DFT calculations to understand the structural stability for sodium extracted phases. From these results, the detailed mechanisms on the loss of reversible capacity are comprehensively discussed.

(EXAFS) spectra were converted from energy to wave vector k and then weighted by k3. Morphology of the samples was observed by scanning electron microscopy (S-5000H, Hitachi) operated at 10 kV of acceleration voltage. Beaker-type and coin-type cells were used to evaluate electrode performance in Na cells. The positive electrode consisted of the active material:acetylene black:PVdF = 8:1:1 in weight ratio. The electrode components were thoroughly mixed in N-methylpyrrolidone to form uniform black slurry, and then the slurry was pasted to stainless steel mesh or on aluminum sheet as current collector and dried in vacuum at 80 °C. A thin piece of sodium metal was used as a negative electrode. The electrolyte solution used was 1.0 mol dm−3 NaClO4 dissolved in propylene carbonate (PC) for beaker- and coin-type cells and 1.0 mol dm−3 NaPF6 dissolved in propylene carbonate for coin-type cells. The electrodes and electrochemical cells were prepared under glove boxes in order to suppress deterioration by water contamination. Electrochemical charge and discharge tests were carried out at room temperature (ca. 25 °C). The structural changes during the charge/discharge tests were examined by an ex-situ XRD technique. The Vienna ab initio simulation package (VASP)42,43 was utilized with the generalized gradient approximation (GGAPBE) + U44,45 and projector-augmented wave (PAW) methods.46 Besides, the on-site Coulomb correction (GGA + U) was included for localized electronic states, and the U value was chosen to be 3.5 eV for Cr 3d states according to the literature.47

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RESULTS AND DISCUSSION Crystal structure of O3-type NaCrO2 was examined by X-ray diffraction (XRD). Figure 1 and Table 1 show an X-ray

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EXPERIMENTAL SECTION NaCrO2 powders were prepared by a conventional solid-state reaction from stoichiometric ratio of Cr2O3 (Kanto Chem. Co., Inc., purity: 98.5%) and Na2CO3 (Kanto Chem. Co., Inc., purity: 99.5%) as starting materials. Mixture of the starting materials is pelletized and heated at 900 °C in Ar for 5 h.33 Crystal structures of the samples were examined by using Xray diffractometer (MultiFlex, Rigaku Co., Ltd.) with Cu Kα radiation source without air exposure by use of a laboratory made attachment of sample holder. The X-ray diffraction (XRD) data were analyzed and simulated by using the Rietveld refinement program RIETAN-FP.40 X-ray absorption spectroscopy (XAS) for the samples was conducted at beamline BL-9C and 12C of the Photon Factory Synchrotron Source in Japan. XAS spectra were collected with a silicon monochromator in a transmission mode. The incident and transmitted X-ray intensity was measured using an ionization chamber at room temperature. The samples were put in a water-resistant polymer film and sealed in the Ar-filled glovebox to minimize influence from moisture during measurement. Analysis of the XAS spectra was carried out using the ATHENA software package based on the IFEFFIT.41 The XAS spectra were normalized with reference samples. The normalized extended X-ray absorption fine structure

Figure 1. Results of the Rietveld refinement on a powder X-ray diffraction pattern of NaCrO2. A schematic illustration of the crystal structure for NaCrO2 drawn using the program VESTA62 is also shown (inset: oxide ions with different positions in a−b planes are denoted as A, B, and C).

diffraction pattern of NaCrO2 synthesized and its refinement results by the Rietveld method. All the diffraction lines were indexed with a space group R3m ̅ as O3-type layered structure, and no diffraction peaks from impurities were observed. A schematic illustration of the O3-type structural model is also shown in Figure 1. Wyckoff positions for each element with B

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The Journal of Physical Chemistry C Table 1. Structural Parameters of NaCrO2 Refined by the Rietveld Method; Isotropic Atomic Displacement Parameter (B) and Occupancy (g) of All Sites Were Fixed formula NaCrO2 space group R3̅m ahex = 2.97610(4) Å, chex = 15.96723(19) Å Rwp = 9.51%, RB = 6.31%, RF = 4.62%, S = 1.20

a

atom

site

x

y

z

Ba/Å2

ga

Na Cr O

3b 3a 6c

0 0 0

0 0 0

1/2 0 0.26406(8)

1.0 0.5 0.5

1.0 1.0 1.0

Not refined.

space group R3̅m used for the refinement were as follows: Na at 3b (0, 0, 1/2); Cr at 3a (0, 0, 0); and O at 6c (0, 0, z) with z ≈ 0.26. From the structural refinement, lattice parameters are determined to be ahex = 2.97610(4) Å and chex = 15.96723(19) Å in a hexagonal setting. These values are identical with those in the previous report.24 The refinement by the Rietveld analysis showed reasonably small R factors (Rwp = 9.51% and S = 1.20) without refinement of a site occupation of chromium ions at 3b site in the interslab layer, suggesting the intermixing between the sodium and chromium ions is negligible. From these results, stoichiometry of NaCrO2 can be regarded as an ideal layered O3-type structure without an antisite defect between sodium and chromium ions. The morphology of NaCrO2 was examined by scanning electron microscopy (SEM). From the SEM images collected at different magnifications shown in Figure 2, the primary particle

Figure 3. (a) Galvanostatic charge and discharge curves for Na// NaCrO2 cells with 1 M NaPF6/PC solution at C/20 rate (12.5 mA g−1) in the voltage ranges of 2.5−3.6 V. (b) Discharge capacity retention of the Na//NaCrO2 cells at C/20 rate (12.5 mA g−1) in the voltage ranges of 2.5−3.6 V.

irreversible capacity at each cycle and capacity fading during extensive cycles would be caused by side reactions such as electrolyte decomposition and/or oxidation of the decomposed compounds produced on surface of counter Na electrode.48 The side reactions would be suppressed by a carbon-coating method to improve the capacity retention.49 The electrochemical reversibility of Na extraction/insertion processes was investigated with different upper cutoff voltages of ca. 3.7, 3.8, and 4.5 V and different Na extraction ranges of x = 0.0−0.5, 0.0−0.7, and 0.0−ca. 0.94, respectively, in Na1−xCrO2 at 1/20 C (12.5 mAg−1). As shown in Figure 4, Figure 2. Morphology of as-prepared NaCrO2 observed by a scanning electron microscope.

size of as-prepared NaCrO2 ranges from approximately 1 to 5 μm, and the primary particles seem to be a single crystallite. No aggregated particles are observed in Figure 2. The highly crystallized NaCrO2 with well-dispersed particles is obtained by a solid-state reaction. The electrode performance of NaCrO2 was examined at room temperature using coin-type cells. Galvanostatic charge/ discharge curves of a Na//NaCrO2 cell in a voltage range of 2.5−3.6 V are shown in Figure 3a. The reversible capacity of 110 mAh g−1 was delivered with a small irreversible capacity, indicating reversible sodium extraction/insertion from/into O3-type NaxCrO2 in the voltage range of 2.5−3.6 V. The cell also exhibited relatively good capacity retention within this voltage range as shown in Figure 3b. When we tested at higher current density of 1C rate or with different electrolyte solution, the cell exhibited almost the same discharge capacity and cycle retention (see Supporting Information, Figure S1). The small

Figure 4. Initial charge and discharge curves of Na//NaCrO2 cells at a rate of 1/20 C (12.5 mA g−1) in the ranges of 0.0 ≤ x ≤ 0.5 and 0.0 ≤ x ≤ 0.7 in Na1−xCrO2 and in a voltage range of 2.5−4.5 V. C

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Figure 5. (left) Ex-situ X-ray diffraction patterns of the Na1−xCrO2 composite electrodes on a stainless steel current collector, which were electrochemically prepared in the sodium cells. (right) Highlighted XRD patterns. Asterisks show stainless steel mesh used as a current collector.

an initial discharge capacity of 118 mAh g−1 was delivered by 0.5 mol of Na extraction with a small irreversible capacity of 8 mAh g−1. When the cell was charged to middle of the plateau region at ca. 3.8 V, the reversible discharge voltage plateau was not seen around 3.8 V. In addition, further charge beyond the 3.8 V plateau region resulted in the significant decrease in the discharge capacity of only 9 mAh g−1 even though almost theoretical charge capacity (∼250 mA g−1) was obtained in charging up to 4.5 V (∼235 mA g−1). The dependency on the upper cutoff voltage is the consistent results with those in our previous report.33 Similar behavior, the dependency of cutoff voltage, is also reported in O3-NaTiO2,21 O3-NaVO2,50 and O3-NaFeO2,21,26,50 and the reversible capacity of these layered materials also drastically decays by charge to higher cutoff voltages. The drastic capacity decays would be attributed to an irreversible migration of transition metals from transition metal oxide slabs to the interslab space. To the best of our knowledge, although the migration of chromium ions from CrO2 slabs to interslab space is reported in Li removal from LiCrO2,51−54 no report is found in the Na analogue system. To understand the capacity decay mechanism for the NaCrO2 electrode, the structural changes were examined by ex-situ XRD measurements. The samples at various states of charges (various sodium compositions) were electrochemically prepared by galvanostatic charging at 1/25 C (10.0 mAg−1) for NaCrO2 composite electrodes. After the charge in the Na cells, the cells were disassembled and the composite electrodes were set in a sample holder without air exposure. The ex-situ XRD patterns of the composite electrodes with different Na amount of x in Na1−xCrO2 are shown in Figure 5. The peaks from a stainless steel mesh used as a current collector in the Na cells are marked with asterisks. Lattice parameters were calculated by Rietveld analysis for the diffraction patterns, and variations in the neighboring Cr−Cr and interslab distances are plotted in Figure 6. In the XRD pattern of as-prepared electrode, x = 0 in Na1−xCrO2, the diffraction lines of the pristine O3 phase are observed. When the sodium ions are extracted to x = 0.09−0.17 in Na1−xCrO2, the 003hex diffraction line located at 17° (2θ) splits into two peaks as show in Figure 5, and thus the O3 phase coexists with a new layered phase. The new layered phase was assigned to a monoclinic O′3 phase with space group C2/ m, which has a distorted lattice compared to an ideal hexagonal

Figure 6. Change in the crystallographic parameters and phase evolution of the Na1−xCrO2 as a function of Na contents x-value; averaged Cr−Cr interatomic distances (bottom) and interslab distances (middle). A galvanostatic charge curve of the Na// NaCrO2 cell is shown for comparison (top).

O3-type unit cell (the prime symbol refers to the distorted lattice). The lattice distortion results in peak separations of the Bragg diffractions, e.g., the 104hex diffraction line in the O3 phase split into 202̅mon and 111mon in the O′3 phase. In the further extraction of sodium ions to x = 0.26, the monoclinic O′3-type layered phase partially transforms into a monoclinic P′3-type layered phase because the intensity of 202̅mon and 111mon lines decreases, while those of 112m ̅ on and 201mon lines obviously increase. The P′3-type structure with space group C2/m also has a distorted lattice of an ideal P3-type structure where sodium ions are located at triangular prismatic sites in the interslab space. For the samples of x = 0.35, 0.44, and 0.52, the single P′3-type phase was observed. The phase transitions from O3- to P′3-type through the O′3 phase during charge in D

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The Journal of Physical Chemistry C the Na amount of 0.0 ≤ x ≤ 0.5 were also reported in the literature.31,34 Since the electrostatic repulsion between facing oxide anions of the MeO2 slabs is increased by the removal of Na ions from the interslab space, the calculated interslab distance based on the lattice parameter c (chex/3 or cmon sin β for the hexagonal or monoclinic lattice, respectively) increased with extraction of Na ions in the range of 0.0 ≤ x ≤ 0.5 in Na1−xCrO2 as shown in Figure 6. On the other hand, Cr−Cr distance (corresponding to the lattice parameter ahex and bmon) shortened, which would be contributed by Na extraction from the interslab space and the oxidation of chromium ions with the change in ionic radii from 0.615 (= Cr3+) to 0.55 (= Cr4+) as discussed later with XAFS data. In addition, the drastic stepwise change in voltage from 3.3 to 3.65 V at x ≈ 0.44 is reversibly seen in charge/discharge curves as shown in Figure 4. Although a Na/vacancy ordering in the interslab space generally brings drastic voltage changes as reported for P3-type NaxVO2,55 clear evidence of sodium/vacancy ordering was not evidenced by using a lab level X-ray diffractometer. Regarding this point, an additional structural study by using synchrotron ex-situ XRD56 successfully observed a diffraction for superlattice originating from the in-plane Na/vacancy ordering in Na0.5CrO2, and the detailed structural studies will be reported elsewhere. By further Na removal to x = 0.87, another O3-type layered phase (O3′ phase) with trace of unreacted P′3-type phase was observed. No significant change is seen in the lattice parameter of ahex because of similar Cr−Cr distance at x = 0.52 and 0.87; however, 003 lines drastically shifted to higher angle, reflecting the shrinkage of the lattice along the c-axis. Interestingly, the intensity of the 003 diffraction line is much smaller compared to those of 104 and 101 lines, suggesting a chromium migration from slabs to interslabs. Such behavior has been also reported in O3-NaTiO2 and O3-NaFeO2.21,26 The crystal structure of O3-NaTiO2 changes into an O′3-type layered structure accompanied by Na extraction, and further Na extraction leads to an irreversible phase transition with lowered intensity of 003 diffraction line.21 The decrease in the intensity of 003 line associated with the irreversible structural change during charge was also observed for the O3−Na1−x FeO2 in x > 0.5.26 Although reduction of intensity for the 003 line and the irreversible structural change are similarly observed in O3NaCrO2, O3-NaCrO2 electrodes undergo a different phase transition compared with those of O3−Na1−xTiO2 and O3− Na1−xFeO2. The ex-situ XRD measurements for the O3− Na1−xCrO2 electrodes in fully charged stage of 4.5 V were, therefore, carried out to investigate detailed mechanisms for the migration of transition metal in NaCrO2. Figure 7 compares the XRD patterns of Na1−xCrO2 electrode after charge to x = 0.7 to 4.5 V and simulated patterns using the models with different degrees for Cr migration. When the cell was charged to x = 0.7, that is the middle position of the charge plateau at ca. 3.8 V, main diffraction peaks in the XRD pattern can be indexed as P′3-type phase without a significant change in lattice volume compared to that of x = 0.52 shown in Figure 5. Asymmetric peak profile is, however, observed for 00l lines, and the structure gradually and irreversibly transforms at the plateau region, which presumably originates from the irreversible capacity degradation after charge beyond x = 0.5 in Na1−xCrO2 as shown in Figure 4. On the other hand, after charging to 4.5 V, where the Na extraction amount x is estimated to be 0.94 according to the state of charge with the assumption that all current passed through the electrode was consumed by sodium extraction process without side reactions,

Figure 7. (a) Ex-situ X-ray diffraction pattern of the Na1−xCrO2 composite electrode after charge to x = 0.7 and 4.5 V in Na coin cell and the simulation patters of Na0.06CrO2 with a Cr migration to tetrahedral and octahedral sites in Na layers. Asterisks show aluminum metal used as a current collector. (b) Crystal structure and interstitial tetrahedral and octahedral sites.

the diffraction lines of Na1−xCrO2 (x = 0.94) can be assigned to a rhombohedral O3-type phase with a space group, R3̅m. The simulated pattern of an ideal layered O3-type Na0.06(Croct)O2 model without chromium migration is, however, clearly different from the observed pattern in terms of the intensity of diffraction lines. Therefore, several models with different chromium migration manners: chromium migration at tetrahedral and/or octahedral sites in the interslab space as represented in Figure 7. The site for tetrahedral chromium ions in the interslabs space was set at 6c (0, 0, z) with z = 1/8. The tetrahedral sites in CrO2 slabs share faces with the neighboring octahedral sites in the slabs, which is expected to induce a very strong repulsive interaction. Therefore, it is anticipated that chromium ions can occupy either of the octahedral and tetrahedral sites in the CrO2 slabs. On the other hand, since Na+ ions at the edge-shared octahedral sites are almost extracted, face-shared tetrahedral sites with three octahedral ones in interslab space would be preferred for chromium ions and octahedral sites in the interslab space can be also occupied in partial by chromium ions. The simulated pattern of Na0.06Croct.0.3(Croct.0.7)O2, in which 30% chromium ions are assumed to migrate into only octahedral sites in the interslab, shows clear difference in relative intensities of 003, 104, and 108/101 lines in comparison to those of the observed pattern. In contrast, when we simulate the XRD pattern for Na0.06Croct.0.1Crtet.0.2(Croct.0.7)O2, in which chromium ions are located at both octahedral and tetrahedral sites in the interslab space, we obtain a quite similar pattern to the observed XRD profile compared with the models with the assumption of chromium migration only to tetrahedral or octahedral sites in the interslab space as seen in Figure 7. These suggest that the E

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charge. The energy shifts of Cr K-edge XANES spectra for pristine, pristine, and half-charged Na1−xCrO2 are all consistent with the reported literature.34 For the electrode sample cycled to 2.5 V (after being charged up to 3.6 V), the spectra return to almost the same profile at the same energy with those of the pristine, indicating reversible chromium redox during charge and discharge. In addition, Cr−O and Cr−Cr bond lengths also change reversibly from the observation of the extended X-ray absorption fine structure (EXAFS) spectra for NaCrO2 electrodes as shown in Figure 8b. Peaks at 1.65 and 2.65 Å (without the calibration of phase shift) are assigned to Cr−O and Cr−Cr coordination shells, respectively. The both Cr−O and Cr−Cr bond length decreased by charge to 3.6 V and increased by discharge to 2.5 V (see Supporting Information in Table S1). This is attributed to a change in ionic radius of chromium ions due to redox of chromium ions, which is well agreed with those of XANES spectra and the electrochemical reversibility in this voltage range. Figure 8 confirms that peak intensity of the white line at 6010 eV decreased with clear change in the profile, and peak intensity at 1.55 and 2.55 Å in EXAFS spectra also drastically decreased for the sample electrodes after being charged to 4.5 V (beyond x = 0.5). Furthermore, the sample discharged to 2.5 V following being charged to 4.5 V showed almost the same XANES and EXAFS spectra of the charged state as expected from the small reversible capacity. Additionally, although the pre-edge peaks due to electric quadrupole transition from 1s to 3d were observed for all samples with some intensity difference, the electrode samples after being charged to 4.5 V has the significantly strong pre-edge peak. K2CrO4 also has the strong intensity of the pre-edge peak at 5992 eV due to electric dipole transition from 1s to p−d hybrid orbital with tetrahedral CrO4 coordination.57 An increase in the pre-edge peak intensity was also reported to be observed for LixMnyCrzO2 and Li[Li0.2CrxCo0.4−xMn0.4]O2 during lithium deintercalation.51−53 In the Li counterpart, chromium ions are thought to more easily migrate to the tetrahedral sites by Li extraction as we described previously.24 In Na1−xCrO2 (0.0 ≤ x ≤ 0.5), during the charge, Cr3+(d3) with 6-fold coordination is oxidized to Cr4+(d2) and the ionic radius changes from 0.615 to 0.55 Å as described for ex-situ XRD above.58 The appearance of strong pre-edge peak after charge to 4.5 V (beyond x = 0.5) also supports the irreversible structural changes and a partial migration of chromium ions from CrO2 slabs to tetrahedral sites in interslab space. Magnetic property data reported by Miyazaki et al. also suggested the existence of Cr4+ in Na0.52CrO2 obtained from NaCrO2 by chemical oxidization with bromine,59 supporting the stability of Cr4+ in the range of 0.0 ≤ x ≤ 0.5 in Na1−xCrO2 unlike LiCrO2-related materials in Li cells. From Cr 2p3/2 XPS spectra of the electrodes before and after charge to 3.6 V, Cr3+ and Cr4+ are evidenced (shown in Figure S2 of the Supporting Information). Cr3+ is oxidized to Cr6+ accompanied by the migration of chromium ions form octahedral to tetrahedral sites in the Li system.52 Ex-situ XRD and XAS studies have clearly revealed that the stable electronic structures and charge compensation mechanisms of NaCrO2 electrode during charge and discharge in the Na cell would be quite different from those of LiCrO2 electrode in the Li cell. Therefore, first-principles DFT calculations were performed to compare the structural stability of O3- and P3type phases for Na1−xCrO2 during charge and further understand the reaction mechanism. It is noted that we do not distinguish P3- and P′3- or O3- and O′3-type structures

migration of chromium ions from octahedral sites in CrO2 slabs to tetrahedral and octahedral sites in the interslab space occurs, which is induced by removal of sodium ions beyond x = 0.5 in Na1−xCrO2. Such migration of chromium ions is irreversible process and therefore disturbs reinsertion of sodium, resulting in the loss of reversibility as the electrode material as shown in Figure 4. Consequently, the reversible range as the electrode material should be limited to 0.0 ≤ x ≤ 0.5 in Na1−xCrO2. X-ray absorption spectroscopy at Cr K-edge was also conducted to examine the change in local structure and oxidation state of chromium ions in Na1−xCrO2 during charge and discharge in the different upper cutoff voltage ranges. Figure 8a shows the X-ray absorption near-edge structure (XANES) spectra of the electrode samples with Cr2O3 and K2CrO4 as references for Cr3+ and Cr6+, respectively. White line at 6008 eV for NaCrO2 clearly shifted to 6010 eV after charging to 3.6 V, suggesting an oxidation of chromium ions during

Figure 8. (a) X-ray absorption near-edge structure (XANES) spectra at Cr K-edges of Na1−xCrO2 and (b) extended X-ray absorption fine structure (EXAFS) spectra of Na1−xCrO2 at Cr K-edges; as-prepared, charged to 3.6 V (3.6 V), discharged to 2.5 V after charge to 3.6 V (2.5 V after 3.6 V), charged to 4.5 V (4.5 V), and discharged to 2.5 V after charge to 4.5 V (2.5 V after 4.5 V) Na1−xCrO2. F

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experimental charge curve in the x ≤ 0.7 region. In contrast, the calculated potential is not in good agreement with the potential of observed plateau in 0.7 ≤ x region. From the ex-situ XRD results, the structure of O3-type phase in the 0.7 ≤ x region is found to be different from that of ideal O3-CrO2 and a migration of chromium ions from the CrO2 slabs to interslab space occurs. Here, to compare the phase stability, we demonstrated the calculations of Frenkel-like defect formation energies for O3-CrO2, in which interstitial Cr ions occupied at tetrahedral or octahedral vacancy sites in interslab layer (Cr migration model). Calculated defect formation energies are listed in Table 2, and both defect formation energies are

due to small cell size in the computations. Using total energies for fully intercalated (x = 0) and deintercalated (x = 1) compositions as reference, formation energies, defined by following equation, were calculated by first-principles calculations. ΔEf (Na1 − xCrO2 ) = E(Na1 − xCrO2 ) − xE(CrO2 ) − (1 − x)E(NaCrO2 )

(1)

Note that the Na/vacancy arrangement in both O3- and P3Na1−xCrO2 was determined as follows; total energies for different 120 arrangements are calculated by ab initio DFT and formation energy of Na-extracted phases are determined as plotted in Figure 9. As shown in Figure 9a, O3-type structure is

Table 2. Relative Formation Energies of O3-Type CrO2 with a Migration of Cr Ion

more stable than P3-type structure for Na1−xCrO2 in 0.0 ≤ x ≤ 0.3 region. On the other hand, in 0.3 ≤ x region, formation energies for O3- and P3-type structure become closer at 0 K, indicating that P3-type structure can be more stable phase at ambient temperature, which is consistent with ex-situ XRD results described above. The redox potential for the sodium removal was calculated by the equation60 E(Na nCrO2 ) − E(Na n − xCrO2 ) − xE(Na) xe

relative energy (eV)

ideal O3-type CrO2 layers octahedral site in Na layer tetrahedral site in Na layer ideal P3-type CrO2 layers

−0.054 −0.011 −0.052

negative, indicating Cr ions spontaneously migrate to the interslab space. Since energy difference between tetrahedral and octahedral interstitial site occupation is rather small, ∼40 meV, both the interstitial sites may be occupied almost evenly by Cr ions due to entropic effect. Chromium ions, therefore, possibly migrate from the octahedral sites in CrO2 slab to tetrahedral sites in interslab space and then to octahedral sites face-shared with the tetrahedral sites by removal of Na+ ions, which is consistent with the discussion of the ex-situ XRD pattern for Na0.06CrO2 and its simulations as shown in Figure 7. All the experimental results and calculations for deintercalated phases lead us to conclude that chromium ions irreversibly migrate from CrO2 slab to interslab space in partial by further extraction from Na0.5CrO2. Models of the chromium migration proposed are summarized in Figure 10. O3-type NaCrO2 transforms to a P′3-type phase through the O′3-type phase by extraction of Na ions in the range of 0.0 ≤ x ≤ 0.5 in Na1−xCrO2. By further extraction of Na ions from the structural lattice in the 0.7 ≤ x region, the P′3-type phase changes to another O3-type (O3′-type) phase by gliding of CrO2 slabs because of its phase instability. Although the P′3-type structure does not have tetrahedral sites in the interslab space, the phase transition yields the vacant tetrahedral sites in interslab space in the O3-type structure. Chromium ions in the octahedral sites in slabs partly migrate to the tetrahedral sites in interslab space and then would further move to octahedral sites from the faceshared tetrahedral sites. The phase transition to the O3′-type phase with chromium migration is irreversible and leads to disturbing insertion of Na ions into the oxide, resulting in a small discharge capacity with a quite large irreversible capacity. The suppression of the structural change into O3′-type phase is expected to be important to obtain a reversible capacity with good cycle retention. Phase transition from O3-type to P3-type phases is not observed for NaFeO2, NaTiO2, and NaVO2 materials, and in general the irreversible migration of transition metals is also observed in these materials. Stabilization of P3type structure during charge is the key factor to extend the reversible range of Na insertion/extraction into/from O3-type layered materials. From the DFT calculations for the phase stability of P3-type structure, 61 substitution of cobalt, manganese, and nickel elements would stabilize the P3-type

Figure 9. (a) Formation energies for the most stable Na/vacancy arrangement as a function of x in O3 and P3 type Na1−xCrO2. (b) Comparison between the experimental voltage profile (black) and calculated averaged voltages of O3-type (blue) and P3-type (red) NaxCrO2.

V=−

site of a migrated Cr

(2)

where E is energy per formula unit at 0 K and e is the elementary charge. Horizontal blue and red lines in Figure 9b show the redox potential for O3- and P3-type phases, respectively. Considering the phase transition from O3- to P′3-type phase at approximately x = 0.3 in ex-situ XRD patterns, calculated redox potential is very close to the G

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Figure 10. A proposed mechanism of chromium migration process on the sodium extraction process.

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ACKNOWLEDGMENTS This study was in part granted by the Japan Society for the Promotion of Science (JSPS) through the “Funding for NEXT Program”, initiated by the Council for Science and Technology Policy (CSTP) and by Ministry of Education Culture, Sports, Science and Technology, Japan (MEXT) program “Elements Strategy Initiative to Form Core Research Center” (since 2012).

structures and extend the reversible limit with a wide compositional ranges as electrode materials for rechargeable Na batteries.38,39

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CONCLUSIONS O3-type NaCrO2 is synthesized by a conventional solid-state reaction, and the influence of different upper cutoff voltage on the crystal structure and electrochemical properties is examined in Na cells. Initial reversible capacity of 110 mAh g−1 is delivered at 1/20C rate in a voltage range of 2.5−3.6 V. Ex-situ X-ray diffraction for NaCrO2 electrodes indicated the structural evolution from O3- to P′3-type phase through O′3-type phase by gliding of CrO2 slabs accompanied by extraction of Na ions, which is consistent result with first-principles calculations. X-ray absorption analysis also reveals the reversible change in electronic state of chromium ions without the formation of Cr6+ at a tetrahedral site during charge and discharge in the range of 0.0 ≤ x ≤ 0.5 in Na1−xCrO2. On the other hand, initial discharge capacity of only 9 mAh g−1 is obtained by charge of 235 mAh g−1 up to 4.8 V. Ex-situ X-ray diffraction and absorption analysis indicates that further extraction of Na ions beyond the plateau region at 3.8 V involves a drastic shrinkage of interslab space accompanied by an irreversible phase transition from P′3- to O3′-type phase. This process is also accompanied by the migration of chromium ions from the slabs to interslab space, resulting in low discharge capacity. Firstprinciples calculations also support the phase stability of the migration model where chromium ion accommodates at both tetrahedral and octahedral sites. By charging and discharging in the range of 0.0 ≤ x ≤ 0.5 in Na1−xCrO2, the irreversible phase transition is suppressed and the reversible capacity with good capacity retention is obtained.

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ASSOCIATED CONTENT

S Supporting Information *

Structural parameters of Na1−xCrO2 obtained from EAXFS measurements, discharge capacity retention of Na//NaCrO2 cells with different electrolyte solution, and Cr 2p3/2 XPS spectra of NaCrO2 electrodes. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] (S.K.). Notes

The authors declare no competing financial interest. H

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