Vacancy Ordering in P2-Type Sodium

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On Disrupting the Na-ion/Vacancy Ordering in P2-Type SodiumManganese-Nickel Oxide Cathodes for Na+-ion Batteries Arturo Gutierrez, Wesley M. Dose, Olaf J. Borkiewicz, Fangmin Guo, Maxim Avdeev, Soojeong Kim, Timothy T. Fister, Yang Ren, Javier Bareño, and Christopher S Johnson J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.8b05537 • Publication Date (Web): 06 Sep 2018 Downloaded from http://pubs.acs.org on September 6, 2018

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The Journal of Physical Chemistry

On Disrupting the Na+-ion/Vacancy Ordering in P2-type SodiumManganese-Nickel Oxide Cathodes for Na+-ion Batteries Arturo Gutierrez,a Wesley M. Dose,a Olaf Borkiewicz,b Fangmin Guo,b Maxim Avdeev,c Soojeong Kim,d Timothy T. Fister,d Yang Ren,b Javier Bareño,a and Christopher S. Johnsona* a

Electrochemical Energy Storage Department, Chemical Sciences and Engineering Division, Argonne National Laboratory, 9700 S. Cass Avenue, Argonne, IL 60439, USA b

X-ray Science Division, Advanced Photon Source, Argonne National Laboratory, 9700 S. Cass Avenue, Argonne, IL 60437, USA c

Australian Centre for Neutron Scattering, Australian Nuclear Science and Technology Organisation, Locked Bag 2001, Kirrawee DC, NSW 2232, Australia d

Molecular Scale Science Department, Chemical Sciences and Engineering Division, Argonne National Laboratory, 9700 S. Cass Avenue, Argonne, IL 60437, USA *

Corresponding Author:

E-mails: [email protected] (C. Johnson)

Abstract An investigation of the electrochemical and structural properties of layered P2Na0.62Mn0.75Ni0.25O2 is presented. The effect of changing the Mn:Ni ratio (3:1) from what is found in Na0.67Mn0.67Ni0.33O2 (2:1) and consequently the introduction of a third metal center (Mn3+) was investigated. X-ray powder diffraction (in-situ and ex-situ) revealed the lack of Na+ion/vacancy ordering at the relevant sodium contents (x = 0.33, 0.5 and 0.67). The Mn3+ in Na0.62Mn0.75Ni0.25O2 introduces defects into the Ni-Mn inter-plane charge order that in turn disrupts the ordering within the Na-plane. The material underwent a P2-O2 and P2-P2’ phase transition at high (4.2 V) and low (~ 1.85 V) voltages, respectively. The material was tested at several different voltage ranges in order to understand the effect of the phase transitions on the capacity retention. Interestingly, the inclusion of both phase transitions demonstrated comparable cycling performance to when both phase transitions were excluded. Lastly, excellent rate performance was demonstrated between 4.3 – 1.5 V with a specific capacity of 120 mAh/g delivered at 500 mA/g current density.

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Introduction Sodium-ion batteries (SIBs) continue to attract a growing portion of the energy storage research community1-3 particularly as possible candidates for large-scale grid storage because sodium is widely available around the globe and is abundant in the earth’s crust. The abundance of sodium reinforces the idea that sodium-ion batteries are a viable way to develop a low cost/high performance alternative to lithium-ion batteries. SIBs function by the same rockingchair principle that lithium-ion batteries do, where the alkali-ion is shuttled back and forth to/from the electrodes through the electrolyte. Presently, the field is working rapidly to improve each of the components of the sodium-ion battery, in particular the cathode. Layered oxide materials have been heavily scrutinized for use as cathodes in lithium-ion batteries4-11 and are currently the most investigated cathode material for SIBs.3, 2, 12-14 Sodium transition metal layered oxides crystallize into several polymorphs determined by the complex interplay of precursor selection, material chemistry, sodium stoichiometry, and processing conditions.15-17 The general structure of these polymorphs can be viewed as a stack of MO2 slabs, where the transition metals occupy the interstitials between two hexagonal O planes, interspersed by Na. The particular structure of a given polymorph is defined by the coordination of the sodium cation (described by a letter; O = octahedral, P = prismatic, T = tetrahedral) and the number of transition metal layers per unit cell,18 indicated by a number following the Na coordination label. Of particular interest for this work is the P2-type polymorph that contains two MO2 slabs per unit cell and trigonal-prismatic-coordinated Na, with the stacking sequence O(A)-M-O(B)-Na-O(B)-M-O(A)-Na. The open prismatic paths in this structure facilitate fast Na+-ion diffusion. Since sodium-cathode chemistries generally exhibit lower energy densities when compared to their lithium analogues, Na0.67Mn1-xNixO2 layered oxides are attractive because the Ni redox 2


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couple increases the average voltage and consequently the energy density of the system. In their seminal study of P2-Na0.67Mn0.67Ni0.33O2, Lu and Dahn noted that the material undergoes a phase transition from P2 to O2 when the sodium content is < 1/3.19 This phase transition comes about as a consequence of increasing O2- - O2 repulsive interactions that occur as the Na+-ion concentration in the material is depleted at the top of charge (> 4.2 V). The repulsion is effectively alleviated as the structure shears, with oxygen layers gliding perpendicular to the caxis accompanied by the creation of stacking faults.19 Although the P2-O2 phase transition is reversible upon discharge, it has been shown that the capacity retention can be significantly improved by limiting the upper cutoff voltage to circumvent the phase transition.20 In addition to the instability introduced by the P2-O2 phase transition, it has previously been determined that Na+-ion/vacancy ordering that occurs at x = 0.33 and 0.5 in NaxMn0.67Ni0.33O2,20 is detrimental to the Na diffusion. The most common approach to overcome both of these challenges in P2Na0.67Mn0.67Ni0.33O2 has been to substitute cations for Ni and/or Mn, including Ti,21 Cu,22-24 and Mg.25-26 Each of these cations was shown to improve the performance by suppressing the P2-O2 phase transition and/or by mitigating the Na+-ion/vacancy ordering that occurs during cycling. In this work, we vary the Mn:Ni ratio from 2:1 found in Na0.67Mn0.67Ni0.33O2 to 3:1 in Na0.62Mn0.75Ni0.25O2 and focus on the effect with regards to the suppression of the P2-O2 phase transition, Na+-ion/vacancy ordering and the overall electrochemical performance. Experimental Synthesis – The samples were prepared by a solid-state reaction of Na2CO3 and Mn0.75Ni0.25(OH)2 (Na0.62Mn0.75Ni0.25O2) or Mn0.67Ni0.33C2O4•2H2O (Na0.67Mn0.67Ni0.33O2). The precursors were mixed in the appropriate mole ratios and fired at 900 oC until the desired phases were achieved. The Mn0.67Ni0.33C2O4•2H2O precursor was prepared by a co-precipitation method 3



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previously described.27 This method involves using an aqueuous solution of NiSO4•6H2O and MnSO4•H2O (Sigma Aldrich, > 99%) mixed with an aqueous solution of Na2C2O4 (Sigma Aldrich, > 99.5%) stirred continuously for 3 h on a hot-plate at ~ 65

o

C. The

Mn0.67Ni0.33C2O4•2H2O precipitate was filtered and washed with distilled water and dried at 110 o

C overnight. The metal ratios for both compounds of interest in this work were measured by

inductively coupled plasma optical emission spectrometry with a Perkin Optima 7300 Dual view ICP-OES at the Analytical Chemical Laboratory 28 at Argonne National Laboratory. Electrochemistry - Electrodes for cell testing were prepared by mixing 84 wt % active material, 8 wt % carbon conducting agent (Super-P Li, Timcal), and 8 wt % polytetrafluoroethylene (PVDF, Solvay 5130). The laminates were dried in air at 75 oC before calendaring and electrodes were punched. Each composite electrode had an active material mass of ∼5 mg. The electrodes were dried in a vacuum oven at 110 °C overnight before coin cell (CR2032) fabrication. The coin cells consisted of metallic sodium anode, 1 M NaPF6 in 3:7 (by wt.) ethylene carbonate/ethyl-methyl carbonate (EC/EMC) electrolyte, glass fiber separator (GF/F, Whatman) and the cathode of the target phase. The electrochemical data were collected at room temperature using a MACCOR battery cycler. X-ray diffraction – High energy synchrotron XRD measurements were carried out at beamline 11-ID-C at the Advanced Photon Source at Argonne National Laboratory (λ = 0.1173 Å). Cells for in situ synchrotron measurements on Na0.62Mn0.75Ni0.25O2 were constructed using modified 2032 coin cells with 2 or 3 mm diameter holes in the cell casing. The cells were hermetically sealed with Kapton windows. Two cells were built and these were subject to the same electrochemical protocol during the beam time. Approximately 11.5 h into the beam time the first cell failed (possibly an issue with uneven stack pressure) and therefore the data from the

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second cell was used from this time. Scans were collected using an a-Si PerkinElmer 1621 area detector with a 90 s exposure time and 15 min between scans. Neutron powder diffraction (NPD) - Neutron powder-diffraction (NPD) data were collected at room temperature on the high-resolution powder diffractometer Echidna at ANSTO. The sample in the form of ~ 5g of powder was loaded into a cylindrical vanadium can and data were collected using neutrons with the wavelength of 1.6215 Å. X-ray Absorption Near Edge Spectroscopy (XANES) – XANES measurements were performed at the MRCAT bending magnet beamline29 at the Advanced Photon Source, Argonne National Laboratory. X-ray absorption spectra were collected in transmission mode through the Na0.62Mn0.75Ni0.25O2 laminates. Energy was scanned by a double-crystal Si(111) monochromator that was detuned by 50% and the incident and transmitted intensity was measured by gas ionization chambers. A Mn reference foil was measured at the same time with every sample for energy calibration, Mn K-edge set to be 6539 eV. Data analysis was completed using the IFEFFIT package.30-31 Results and Discussion I.

Structure of the Pristine Material

The sodium ions in P2-type structures occupy two distinct sites where they either share faces (Naf) or edges (Nae) with MO6 octahedra.32 The relative occupation of the two sites is governed by the complex interaction between Na+ - Na+ repulsion, Na+ - Mn+ interaction, and transition metal charge ordering. The overall occupancy ratio between the two sodium sites leads to several distinct patterns of in-plane Na+-ion/vacancy ordering: i.) honeycomb33 ii.) diamond iii.) row34 and iv.) large zigzag (LZZ). Lee et al. was first to identify LZZ sodium ordering in NaxMn0.67Ni0.33O2 (x = 0.67) by noting the superstructure peaks in SXRD correspond to a d5



spacing consistent with the average distance between nearest neighbor sodium ions in the LZZ pattern.20 The SXRD collected on Na0.62Mn0.75Ni0.25O2 and Na0.67Mn0.67Ni0.33O2 (nominal compositions – ICP-OES and neutron refined compositions can be found in Supplementary Information Table S1) for this work are shown in Figure 1a. Both samples showed very similar patterns and the majority of the reflections differ only by small shifts in 2θ and by relative peak intensities; differences which are controlled by the underlying chemistry of the two samples. The most significant distinction between the two XRD patterns is the clear LZZ superstructure peaks found at 2.05 and 2.14° for Na0.67Mn0.67Ni0.33O2 (see inset of Figure 1a) corresponding to dspacing values of 3.144 Å and 3.274 Å, respectively, that are absent in the Na0.62Mn0.75Ni0.25O2 pattern.

Figure 1. (a) High energy synchrotron X-ray powder diffraction (SXRD) data for Na0.62Mn0.75Ni0.25 and Na0.67Mn0.67Ni0.33O2 showing evidence of ordering reflections at 2.05 and 2.14° 2θ (see inset) observed in the sample with a Mn:Ni ratio of 2:1 which are absent in 3:1; (b) 1st cycle charge/discharge curves for Na0.62Mn0.75Ni0.25 and Na0.67Mn0.67Ni0.33O2 between 4.3 – 1.5 V (15 mA/g) revealing the Na+-ion/vacancy ordering in the latter observed as steps in the voltage profile.

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Figure 2. (i) Transition metal plane in Na0.67Mn0.67Ni0.33 showing the Mn4+ (magenta) and Ni2+ (green) ordering described by the (√3 x √3)-R30° (black diamond; also called large hexagonal unit cell) of the underlying hexagonal plane. (ii) Schematic of the large hexagonal unit cell showing: three unique transition metal sites (i.e., A, B and C), face-sharing sodium (Naf ; gray) underneath each transition metal and edge-sharing sodium (Nae ; black), and oxygen (iii) oblique view of the large hexagonal unit cell (Nae and oxygen omitted) showing the Mn-Mn (blue box) and Mn-Ni (red box) dumbbells formed with Naf sites (iv) (√3 x √3)-R30° from (i) reconstructed to indicate the two unique Naf sites (Mn-Mn or Mn-Ni) matching the box colors used in (iii). Fluorescent green dots indicate three Nae vacancies surrounding one filled Naf site forming a (V3Naf)2- cluster.

The LZZ superlattice peaks have also been identified in Na0.67Mn0.67Ni0.33-yCuyO2 (0 ≤ y ≤ 0.33).24 The authors of that work noted that the superlattice peaks increased in intensity with the

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amount of copper content. Additionally, the superlattice peaks appear to be present in the work done by Kubota et al. on the same series (i.e. Na0.67Mn0.67Ni0.33-yCuyO2) although not mentioned by the authors.23 Furthermore, the superlattice peaks were present in NaxMn0.65Ni0.20Cu0.15O2 when x = 0.67 and 0.75 but not for x = 0.5.22 Finally, the LZZ superlattice peaks appear to be present in the work done by Mason et al. on Na0.67Mn0.67Cu0.33O2 but were attributed to impurities.35 The suitable chemistry (i.e. sodium content, TM ratios, TM oxidation states, etc.) in P2-type sodium transition metal oxides that allows the LZZ pattern to occur appears to be narrow especially when considering the large number of materials tested in the literature and the seemingly low number of samples where the LZZ superlattice peaks have been observed. The effect of the underlying chemistry to either promote or prevent LZZ ordering is discussed below. The ordering of transition metals plays a significant role in the stabilization of the LZZ Na ordering. Figure 2a shows the ordering of Ni2+ (green) and Mn4+ (magenta) ions in the TM planes of in Na0.67Mn0.67Ni0.33O2. Each Ni2+ is surrounded by six Mn4+ ions, and the plane can be viewed as a network of edge-sharing NiMn6 hexagonal units. This ordering can also be described as a (√3 x √3)-R30° reconstruction of the underlying hexagonal plane, which unit cell (large hexagonal unit cell - black diamond in Figure 2a)15 contains three distinct sites as shown in Figure 2b): i.) Site A where Mn is at the corners of the unit cell; ii.) Site B where Mn is 1/3 of the way along the long diagonal of large hexagonal unit cell; iii.) Site C where Ni is 2/3 of the way along the long diagonal of large hexagonal unit cell. The stacking of O and TM planes in the Na0.67Mn0.67Ni0.33O2 structure follows an O-TM-OO-TM-O pattern; i.e., TM atoms stack directly on top of each other, at the center of MO6 octahedra. Each pair of consecutive O planes forms a network of trigonal prismatic cages. As mentioned briefly above, the Na atoms occupy two kinds of sites within these cages, face sharing 8


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(Naf) or edge sharing (Nae), classified by how sodium polyhedra connect to the nearest-neighbor TM polyhedra. The Nae sites (black) sit at the center of a triangle formed by transition metals (See Figure 2b), while the Naf sites lie directly between two TM in consecutive planes, as illustrated in Figures 2b (top view) and 2c (oblique view, with O and Nae sites omitted). The stacking of TM planes, illustrated in Figure 2c, does not place Ni atoms in consecutive TM planes directly on top of each other,15 but rather form (one) Mn-Mn and (two) Mn-Ni dumb bells. This breaks the simple hexagonal symmetry of the Naf sites, according to whether the Naf site lies between two equivalent (Mn-Mn dumbbell) or inequivalent (Mn-Ni dumbbells) transition metals. These two types of Naf sites are indicated in Figure 2c by a blue and red box, respectively. With this TM plane stacking, there are two Mn-Ni Naf sites per Mn-Mn Naf site, also forming a (√3 x √3)-R30° pattern, as illustrated in Figure 2d. Here, red circles represent MnNi Naf sites and blue circles Mn-Mn Naf sites, matching the box colors used in Figure 2c. In the LZZ-ordered Na0.67Mn0.67Ni0.33O2 structure, the distance between two occupied Naf sites (~ 3.2 Å) confirms that Na atoms occupy the alternating Mn-Ni Naf sites indicated by (large) gray dots in Figure 2d. The occupied Naf sites form a 2x2 reconstruction of the Mn-Ni periodicity in the TM plane, indicated by the black diamond in Figure 2d, and account for 25% of the Na in Na0.67Mn0.67Ni0.33O2. The remaining Na ions (6 per 2x2 cell) occupy the Nae sites (small black) as illustrated in Figure 2d. Simple electrostatics arguments can be employed to understand the LZZ ordering. To begin with, Na vacancies (V) are equivalent to a negative charge in the Na plane, with respect to the average charge in the plane. In the LZZ-ordered-Na0.67Mn0.67Ni0.33O2 structure, three Nae vacancies (fluorescent green in Figure 2d) group around each Naf occupied site (gray dots), effectively forming a (V3Naf)2- cluster and ordered arrangements maximizing the distance 9



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between these (V3Naf)2- clusters should be favored. The 2x2 arrangement of (V3Naf)2- clusters, with two clusters per unit cell, keeps them as far apart as possible, forming zig-zag rows where every other Naf site is occupied, and minimizing their Coulombic repulsion. Furthermore, Lee et al. rightly affirmed that Nae sites were more energetically favored when compared to Naf sites due to the unshielded electrostatic repulsion between Na+ and TMn+ through polyhedron faces. Yet from an electrostatics interpretation, the distinct dumbbells (i.e. M-M and Mn-Ni) created by the stacking sequence also suggests that all Naf sites are not energetically equivalent. One could argue that sodium would preferentially fill Mn4+-Ni2+ Naf sites in order to avoid the extra repulsion felt in the more highly charged Mn4+-Mn4+ Naf sites. The authors of this work, in fact, believe that the Na+-ion/vacancy ordering observed in Na0.67Mn0.67Ni0.33O2 is influenced significantly by the tendency of sodium to avoid the Mn-Mn configuration

and

can,

therefore,

be

altered

by changing

the

Mn:Ni

ratio

(i.e.

Na0.62Mn0.75Ni0.25O2) and introducing defects in the Ni-Mn charge order. The hypothesis that Na+-ion/vacancy ordering in P2-type materials can be prevented by breaking the charge ordering of transition metal layers was previously suggested for Na0.6Cr0.6Ti0.4O2.36 Wang et al. proposed that choosing transition metals with similar ionic radii but different redox potential would alter the ordering of Cr3+ and Ti4+ within the transition metal layer (i.e. intra-layer ordering). Unlike X-ray diffraction, neutron powder diffraction (NPD) is very sensitive to Mn-Ni ordering due to their very strong scattering contrast (coherent scattering lengths being -3.73 fm and 10.3 fm, respectively) and can, therefore, be used to determine if any change in intra-layer ordering occurs between Na0.67Mn0.67Ni0.33O2 and Na0.62Mn0.75Ni0.25O2. Previously, Lu et al. observed additional NPD peaks for Na0.67Mn0.67Ni0.33O2 that could not be indexed by a small (hexagonal) P63/mmc unit cell (i.e., with a single transition metal site per 10


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Figure 3. Experimental NPD data (λ=1.6215Å) for the two samples (symbols) and the pattern simulated for the Mn/Ni-disordered model (solid green line)15. Arrows show peaks that cannot be indexed by a small P63/mmc unit cell and instead are assigned to superstructure formation of Mn-Ni ordering (i.e. Ni surrounded by six Mn).

basal plane and a ~ 3Å, c ~ 11Å).15 The additional peaks were instead assigned to a (√3 x √3)R30° superstructure (space group P63) formed by Mn-Ni intra-layer ordering (i.e. Ni surrounded by Mn as shown in Figure 2a). Similarly, the NPD data collected for this work (Figure 3, full details in Supplementary Information) clearly shows the same superstructure peaks (e.g., 2θ between 20° and 30°) are present in both Na0.67Mn0.67Ni0.33O2 and Na0.62Mn0.75Ni0.25O2. This suggests that intra-layer Mn-Ni ordering does not play a significant role in the loss of the LZZ Na+-ion/vacancy ordering in the Na0.62Mn0.75Ni0.25O2 sample. 11



Instead, the nuance proposed in this work is that charge ordering between transition metal layers (i.e. inter-layer ordering see Figure 2c) facilitates Na+-ion/vacancy ordering to occur in Na0.67Mn0.67Ni0.33O2. To reiterate briefly what was stated earlier, as the Mn:Ni ratio is increased to 3:1 in Na0.62Mn0.75Ni0.25O2, the additional Mn (~ 7 mol.%) ends up on sites previously occupied by Ni. This changes 25% of the Mn-Na-Ni dumbbells found in Na0.67Mn0.67Ni0.33O2 to Mn-Na-Mn dumbells in Na0.62Mn0.75Ni0.25O2. The additional repulsion felt by Na in a Mn-Na-Mn dumbbell, due to the extra charge on Mn compared to Ni (3+ vs. 2+), consequently makes the site less energetically favorable for occupancy, which disrupts the LZZ Na+-ion/vacancy ordering. It follows that increasing the concentration of higher energy Mn-Na-Mn sites in the structure disrupts the LZZ pattern.

Figure 4. In-situ XRD data (left) collected during the first cycle charge/discharge (right) for Na0.62Mn0.75Ni0.25 between 4.3 – 1.5 V (vs. Na+; 25 mA/g) at room temperature. The XRD and electrochemical profile are color-coded to indicate the regions where the P2-O2 (orange) and P2P2’ phase transition occur.

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II.

Structural Evolution During Electrochemical Cycling

The structural evolution of Na0.62Mn0.75Ni0.25O2 was monitored using in-situ SXRD during a galvanostatic (20 mA/g) charge-discharge cycle between 4.3 – 1.5 V (vs. Na+). Figure 4 (left) contains a stack plot of the θ-2θ SXRD patterns acquired at different states of charge during the cycle and (right) the corresponding cell voltage (x-axis) as a function of cycling time (y-axis). The voltage profile and corresponding cell parameters (volume, c and a) are shown in Figure 5.

Figure 5. Initial voltage profile (top) and corresponding evolution of the volume, c and a with respect to time for Na0.62Mn0.75Ni0.25 between 4.3 – 1.5 V (vs. Na+; 25 mA/g) at room temperature. The shaded regions indicate the P2-O2 (orange) and P2-P2’ (green) phase transition. The square and circle data points represent two different cells that underwent the same test. Rietveld refinements were not conducted between 15.5 – 16.5 h where the P2 and P2’ phases were simultaneously present making refinement difficult due to similar lattice parameters.

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During the initial stages of the charging cycle, up to about 2 h (~ 50 mAh/g), the 002 peak shifts to lower and the 100 peak to higher 2θ, reflecting the linear evolution of the a and c lattice parameters (respectively decreasing and increasing, see Figure 5) as sodium is extracted, while the cell voltage increases to ~ 3.70 V. The a-contraction is caused by the decreasing size of the transition metals as they are oxidized during charging; while the c-expansion is a result of the increasing electrostatic repulsion between oxygen layers as sodium is removed from the structure. The combined effect of the changes in a and c is a small, linear reduction of the unit cell volume. Between 2 and 2.5 h, the cell voltage and c lattice parameter remain stationary, while a continues to decrease. Correspondingly, the decrease in unit cell volume accelerates. While the constant value of the cell potential in this range points to a two-phase equilibrium, there are no indications in the XRD patterns as to what the second phase may be; i.e. no additional reflections, nor peak splitting or significant broadening are evident. One possibility is that regions of low (estimated content at t = 2h) and high (estimated content at 2.5 h) Na content can coexist within the P2 host without lattice parameter changes, resulting in a high- and low- Na content phase equilibrium rather than solid-solution behavior. On continued charging, between 2.5 h and 7 h, the cell voltage remains constant at 4.20 V, until it rapidly increases at the end of the charge cycle. During this time, a, c, and the unit cell volume remain constant. The voltage plateau is indicative of a two-phase equilibrium. Indeed, at t ~ 5 h (or ~ 55% into the plateau) a new reflection appears at 2θ = 1.49° (d-spacing = 4.515 Å) in the SXRD patterns, the intensity of which gradually increases during the rest of the charge cycle (see inset of Figure 4 – right). This reflection is characteristic of the O2 NaxMO2 polymorph, indicating that the 4.20 V plateau is due to a P2-O2 phase transformation. We attribute the absence of this characteristic reflection in

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the SXRD patterns acquired in the first half of the 4.20 V plateau to an insufficient fraction of the O2 phase to provide adequate signal. The re-sodiation of the P2 phase starts with a corresponding 4.10 V plateau between t = 7 h and 10 h, corresponding to ~ 75% of the capacity of the 4.20 V charge plateau. During this time, the characteristic O2 reflection at 2θ = 1.49° diminishes gradually, until it disappears at the end of the plateau. As was the case during the corresponding charge, the unit cell parameters remain constant in the discharge plateau. While the small difference between the charge and discharge voltages can be attributed to cell polarization, the ~ 4.20 V plateau and corresponding SXRD patterns are consistent with a partially reversible P2-O2 phase transformation at this potential. At the end of the 4.1 V discharge plateau, at t = 10 h, the cell voltage rapidly decreases to ~ 3.6 V, and remains at this value up to t = 12 h. Mirroring the behavior that occurred during charge, the unit cell volume and a increase slightly in this voltage range, while c decreases slightly. While c follows a continuous trend through this voltage range, both a and the unit cell volume increase abruptly at t ≲ 12 h. Further discharge between 3.5 V and 1.85 V proceeds predominantly via a solid solution reaction, evidenced by a continuous change to the a- and the c-axis (see Figure 5 between 11 – 14 h). As the voltage reaches ~ 1.85 V, a new plateau emerges, indicative of a phase transition between two Na-prismatic phases, P2 and P2’. This phase transformation shows up in the SXRD patterns as two peaks near 2.7°. Due to the subtle nature of the phase transition (i.e. similar lattice parameters), Rietveld refinements were not conducted in the region where the two phases were present (see Figure 5 between 15.5 – 16.5 h). At the end of discharge (1.5 V) the sample stoichiometry is Na0.89Mn0.75Ni0.25O2 and is possibly made up of two phases: orthorhombic P2’

15



(Cmcm) and monoclinic P2’ (C2/n), which are difficult to distinguish by means of XRD experiments.37 To understand the effect that changing the Mn:Ni ratio has had on the structural evolution we compare the in-situ SXRD results collected in this work on Na0.62Mn0.75Ni0.25O2 with recent work on Na0.66Mn0.67Ni0.33O2.38 Firstly, changing the ratio between Mn and Ni has not subdued the P2 to O2 transition when cycled above 4.2 V. In contrast, substitution (Mg38, 26 or Cu23, 25, 24) into the parent material (i.e. Na0.67Mn0.67Ni0.33O2) for either Mn or Ni has proven to be an effective way to circumvent the P2 to O2 phase transition and instead promote a P2 to OP4 phase transition. The OP4 phase is an intermediate between the P2 and O2 phases consisting of alternate stacking of octahedral (O2) and trigonal prismatic (P2) layers along the c-axis and, therefore, also reduces the damage of the structure caused by extensive gliding of layers.39 Secondly, increasing the Mn:Ni ratio from 2:1 to 3:1 is effective in subduing the Na+ion/vacancy ordering. As discussed above, the LZZ ordering at x = 0.67 is disrupted by the additional manganese which changes the M-Na-M dumbbell environment of the Naf site. In NaxMn0.67Ni0.33O2 sodium also orders into rows at concentrations of x = 0.5 and 0.33.20, 19 These are observed as steps or plateaus in the voltage profile between 3.0 – 3.75 V, as seen in Figure 1b. Using in-situ SXRD, distinct phase transformations associated with these sodium ordered intermediate phases have been observed.38 In contrast, altering the Mn:Ni ratio in NaxMn0.75Ni0.25O2 gives rise to a smooth voltage profile (see Figure 1b), and continuous evolution of the lattice parameters (Figure 5) at sodium concentrations of 0.5 and 0.33; i.e., the intermediate sodium ordered phases are not observed. Since the plateaus are undesirable due to the deleterious effect they have on cycle life and rate performance, substitution for either Ni or Mn has been used widely in order to suppress the 16


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biphasic behavior during cycling. For instance, the voltage steps were effectively suppressed by substituting Ti in for Mn in Na0.67Ni0.33Mn0.67-yTiyO2 (0 ≤ y ≤ 0.67). It is noteworthy that the Na+ion/vacancy ordering was disrupted throughout the whole range of y even when the amount of Ti was small. More commonly, the substitution is for Ni as in the case of Cu in Na0.67Mn0.67Ni0.33yCuyO2.

In this case, the voltage steps can be seen in the parent material (i.e. y = 0) but are

subdued (Figure 2 of referenced work) when both Ni and Cu are in the lattice.23 The work done by Zheng et al. on the same series (i.e. Na0.67Mn0.67Ni0.33-yCuyO2) had the same effect, namely when both Cu and Ni are in the lattice the voltage steps are smoothed but remain for the end members in the series (i.e. y = 0 and 0.33).24 Wang et al. carried out a very similar work where Mg was used instead of Cu (i.e. Na0.67Mn0.67Ni0.33-yMgyO2) and observed less apparent steps in the voltage profile especially when y ≥ 0.1. In addition, Singh et al. noted that although the voltage steps were still present in Na0.67Mn0.7Ni0.3-yMgyO2 (y = 0.05 and 0.1), they were suppressed to a greater extent with increased Mg content.25 In each of the cases mentioned thus far where substitution was used, either for Ni or Mn, it seems that having a third metal center (redox active (Cu) or inactive (Mg or Ti)) in the pristine material is an effective way to disrupt the ordering that occurs at sodium contents of 0.5 and 0.33. If only two metals are present in the pristine material, as in the cases of Na0.67Mn0.67Ni0.33O2 (Mn4+ and Ni2+) and Na0.67Mn0.67Cu0.33O2 (Mn4+ and Cu2+) the ordering at sodium contents of 0.5 and 0.33 still occur as in the two examples mentioned above. It should be noted that Mason et al. did not observe any voltage plateaus in their investigation of Na0.67Mn0.67Cu0.33O2.35 It is hard to know with certainty the reason for the lack of voltage plateaus in that work. Several plausible contributing factors, such as synthesis conditions and charge/discharge currents, appear to be very similar in both studies.35, 24 Of significance is the

17



different electrolyte salts used by the two research groups (NaClO4 – Mason et al. ; NaPF6 Zheng et al.) and previous work demonstrating that a change from NaClO4 to NaPF6 resulted in wellresolved voltage plateaus in the charge/discharge curve of Na0.7CoO2 for the latter and not the former.40 To the best of our knowledge a study of Na0.67Mn0.67Mg0.33O2, where only Mn4+ and Mg2+ are present in the pristine material, does not exist. The closest composition studied is Na0.67Mn0.72Mg0.28O2 and it exhibits a very smooth voltage profile.41 If we assume the oxygen content is 2 for this material, then the Mn oxidation state would be 3.84+ (15% - Mn3+ and 85% - Mn4+) and perhaps a sufficient amount of a third metal center (see Supplementary Figure S1 for Mn K-edge data) has been introduced into the lattice to suppress the ordering. In effect, only two metal centers (Mn3+ and Mn4+) are present in Na0.67MnO2, which also exhibits voltage steps attributed to Mn3+/4+ charge ordering and/or Na+-ion/vacancy ordering.42 Yet, the inclusion of relatively small amounts of a third metal center (Mg2+) in Na0.67Mn1-yMgyO2 (y = 0.05, 0.1) is sufficient to modify the biphasic behavior as evidenced by the smooth voltage profile.42 A similar smoothing of the voltage profile was observed when a third metal center was used in Na0.67Mn0.9M0.1O2 (M = Mg, Ti, Co, Ni, Cu and Zn).43 Indeed, the postulate that three metal centers suppresses the biphasic behavior during cycling is applicable to Na0.62Mn0.75Ni0.25O2 studied in this work. This material also has a Mn oxidation state of 3.84+ signifying a small fraction of Mn3+ is present in the lattice and likewise exhibits a smooth voltage profile at sodium contents of 0.5 and 0.33, which is in stark contrast to Na0.67Mn0.67Ni0.33O2 (see Figure 1b). A similar smoothing of the voltage profile was observed when varying the Ni and Mn ratios in MnThe lack of Na+-ion/vacancy ordering at any concentration (i.e. x = 0.33, 0.5 and 0.67) in

18


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NaxMn0.75Ni0.25O2 suggests that the self-introduction of an additional metal center (Mn3+) by simply changing the Mn:Ni ratio is an effective way to modify the electrochemistry. III. Extended Cycling Performance The cycling performance for the material of interest in this work was tested with a couple of different upper and lower cutoff voltages, as shown in Figure 6a-d. The upper cutoff voltages were chosen to isolate the effect of the P2 to O2 phase transition; 4.05 V (before the transition)

Figure 6. Galvanostatic charge/discharge profiles for Na0.62Mn0.75Ni0.25O2 cycled between (a) 4.05 – 2.0 V (b) 4.3 – 2.0 V (c) 4.05 – 1.5 V and (d) 4.3 – 1.5 V. (e) The capacity retention for the voltage ranges tested in (a-d). All cycling done at room temperature at 15 mA/g.

and 4.3 V (after the transition). It has been suggested elsewhere that the P2 to O2 phase transition is most detrimental to the cycle life of the P2 material because of the dramatic changes (i.e. volume expansion and/or gliding) that occur to the oxygen lattice during the transition.21, 20 The lower cutoff voltage was chosen to isolate the effect of the P2 to P2’ (Cmcm and/or C2/n) phase transition; 2.0 V (before) and 1.5 V (after). The lowering of the voltage allows for the

19



introduction of more Jahn-Teller active Mn3+ and may contribute to poor cycling performance. The cycle life of Na0.62Mn0.75Ni0.25O2 for each of these voltage ranges is shown in Figure 6e. The capacity retention for each voltage range was calculated using the 8th cycle capacity (dashed vertical line) in order to include the break-in experienced when using the widest voltage range (4.3 – 1.5 V). The initial capacity was 88 mAh/g (all other capacities are normalized to this capacity when a comparison is made) when cycled between 4.05 – 2.0 V and 89% capacity retention was achieved after 50 cycles. The initial capacity increases to 160 mAh/g when the upper cutoff voltage is increased to 4.3 – 2.0 V. This represents an 80% increase in capacity but at the expense of inferior capacity retention (79%). These results fit well with previous reports where limiting the upper cutoff voltage to exclude the P2-O2 phase transition in P2-type materials improved cycle life.20-21, 24 The capacity also increases (152 mAh/g) when limiting the upper cutoff voltage to exclude the P2-O2 transition while also dropping the lower limit (i.e. 4.05 – 1.5 V), presumably by including the Mn4+/3+ redox couple (see Supplementary Figure S1). This is an increase in capacity of ~ 73% but also exhibits poor capacity retention (73%). The initial capacity is greatest (185 mAh/g) when cycled between 4.3 – 1.5 V, representing a 110% increase in capacity. Surprisingly, the capacity retention exhibited for the wider voltage window (84%), which includes both phase transitions, was better than when the voltage range only allowed for one of the transitions. We found that the Coulombic efficiency for the first cycle that includes the whole voltage window in question (i.e. 2nd cycle) had a positive correlation to the capacity retention behavior. For instance, the Coulombic efficiency for the 2nd cycle (Figure 6b and 6c) was < 90% when the voltage range included either the P2-O2 or P2-P2’ phase transition and resulted in capacity retentions of 79% and 73%, respectively. In contrast, the capacity retention improved to 20


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84% and 89% when the Coulombic efficiency was > 95% for the initial cycles (Figure 6a and 6d).

Figure 7. The 2nd cycle and 50th cycle (insets) dQ/dV of Na0.62Mn0.75Ni0.25O2 for the various voltage windows tested. The voltage windows tested are labeled in each of the figures. All cycling done at room temperature at 15 mA/g. Shaded gray regions in (b) and (c) represent capacity that is only available after sodium has been intercalated below 2.0 V. Shaded red region in (d) represents capacity only accessed when the P2-O2 phase transition is not included.

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Although a correlation is noted between the Coulombic efficiency and capacity retention, the relationship between the two parameters is not always one of causation, as previously noted.44 One possible reason the initial Coulombic efficiency may have improved/worsened when changing the operating voltage window is due to polarization effects. Briefly, a potential shift (i.e. polarization) away from the equilibrium potential occurs when passing a direct current (i.e. Faradaic process) through an electrode. If the potential shift is large enough it could limit the reversible capacity, and in turn the Coulombic efficiency, by not allowing access to available sodium sites. Dramatic shifts in the redox potentials should be observable by dQ/dV if changing the voltage window in turn altered the polarization of the electrode. Figure 7a-d compares the dQ/dV plots for Na0.62Mn0.75Ni0.25O2 tested in different voltage windows. The “baseline” voltage window (4.05 – 2.0 V) is shown in Figure 7a, along with the voltage window that includes the P2-O2 phase transition (4.3 – 2.0 V). There is no apparent polarization difference in the dQ/dV when the P2-O2 phase transition was included by raising the upper cutoff to 4.3 V. Nevertheless, there was a drop in the initial Coulombic efficiency when the upper cutoff voltage was increased (95% vs. 89% see Figure 6a and 6b). The capacity loss is associated with the P2-O2 phase transition as shown in Figure 6b (black arrow). In fact, capacity loss occurred anytime the upper cutoff voltage included the P2-O2 phase transition (i.e. Figure 6b and Figure 6d). Yet, the capacity retention and the initial Coulombic efficiency results were different depending on the lower cutoff voltage chosen (4.3 – 2.0 V vs. 4.3 – 1.5 V). The dQ/dV for these two voltage windows (Figure 7b) also shows little to no change in polarization suggesting it plays an inconsequential role in the change in performance. The most obvious difference between the dQ/dV for these two voltage windows is the extra capacity (shaded region in Figure 7b) accessed when the lower voltage cutoff is 1.5 V. Actually, accessing the sodium

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sites available < 2.0 V allows the recovery of any capacity lost at higher voltages and, therefore, improved the initial Coulombic efficiency. When the upper cutoff voltage was kept at 4.05 V and the lower cutoff varied similar results were observed (Figure 7c), namely no changes in polarization and some available sodium sites are accessed only < 2.0 V. Yet, the performance (i.e. capacity retention and Coulombic efficiency) worsened when lowering the cutoff voltage to 1.5 V in this scenario (4.05 – 2.0 V vs. 4.05 – 1.5 V) but improved when the upper cutoff voltage was 4.3 V (4.3 V – 2.0 V vs. 4.3 V – 1.5 V) Why? To gain some insight we compare the last scenario (see Figure 7d) that considers the effect of excluding/including the P2-O2 phase transition when the P2-P2’ phase transition is also included (4.05 – 1.5 vs. 4.3 – 1.5 V). First we consider the effect of the P2-O2 phase transition on the initial Coulombic efficiency. Interestingly, the reason for a poorer Coulombic efficiency between 4.05 – 1.5 V may be associated with a lack of available sodium sites < 2.0 V. The lack of available sodium sites is discernible by the telltale signs of concentration polarization at low voltage present in Figure 6c (i.e. “knee” voltage) but absent from Figure 6d. This means that any capacity loss that occurred at higher voltage (black arrow in Figure 6c) cannot be recovered on discharge since all of the available sites < 2.0 V are filled, which in turn lowered the Coulombic efficiency. Additionally, the polarization between charge and discharge is greater at lower voltages when the P2-O2 phase transition is included. After 50 cycles the polarization at lower voltages has reversed and is greater for the sample that excluded the P2-O2 phase transition (see inset of Figure 7d). This is in contrast to all other conditions (see insets of Figure 7a-c) where the polarization after 50 cycles appears similar for the voltage windows compared. The larger

23



polarization observed after 50 cycles for 4.05 – 1.5 V correlates well with the poor capacity retention (see Figure 6e). As mentioned above, it has previously been suggested that the large structural changes associated with the phase transitions leads to poor capacity retention. The results presented here suggest that the effect of the phase transitions can be somewhat alleviated by choosing an appropriate lower cutoff voltage that allows for greater utilization of the available sodium sites at low voltage to recover any lost capacity during initial cycles. The full answer is most likely complex and merits further investigation but recent work on Na0.67Mn0.9M0.1O2 supports our conclusion that large distortions, such as those occurring in phase transitions, may not always be detrimental to the overall performance of the cathode material.43

Figure 8. Discharge rate capability of Na0.62Mn0.75Ni0.25O2 between 4.3 – 1.5 V (vs. Na+). The cell was charged to 4.3 V at 15 mA/g followed by a discharge at different rates between 1 – 500 mA/g.

The rate capability of Na0.62Mn0.75Ni0.25O2 was examined between 4.3 – 1.5 V (vs. Na+) and the results are shown in Figure 8. The electrode demonstrated excellent rate performance

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delivering 224 mAh/g at 1mA/g and retaining 54% (120 mAh/g) of the capacity at 500 mA/g (> 2C).

Similar results have been shown using Mg, Cu or Ti substitution. For example,

Na0.67Mn0.7Ni0.2Mg0.1O2 delivered ~ 120 mAh/g at 400 mA/g,25 Na0.67Mn0.50Ni0.33Ti0.166O2 delivered ~ 90 mAh/g at 484 mA/g,21 and Na0.67Mn0.67Ni0.25Cu0.08O2 delivered ~ 100 mAh/g at 340 mA/g. Still, a direct comparison between these results is difficult because different voltage windows were used when tested. Regardless, the excellent rate capability demonstrated by Na0.62Mn0.75Ni0.25O2 between 4.3 – 1.5 V is somewhat at odds with previous reports for Na0.67Mn0.67Ni0.33O2 where it was proposed that the P2-O2 phase transition, that occurs above 4.1 V, is detrimental to the rate performance of the material. A more detailed study to understand the underlying reason for the improved rate performance is ongoing. 4. Conclusions P2-type sodium-manganese-nickel layered oxides are a low cost/high performance cathode material that offer real promise for use in large-scale applications such as grid storage. The Na0.67Mn0.67Ni0.33O2 cathode material is among the most promising within the P2 family but suffers rapid capacity fade and undergoes biphasic reactions during cycling. Substitution of a third metal center into Na0.67Mn0.67Ni0.33O2, such as Mg, Cu and Ti, has been effective in subduing the Na+-ion/vacancy ordering and suppressing the P2-O2 phase transition. The approach used in this work was to bypass the need for substitution and instead self-introduce a third metal center (Mn3+) into the lattice by varying the Mn to Ni ratio. Although this approach did not subdue the P2-O2 phase transition it did prove to be an effective way to subdue the Na+ion/vacancy ordering at all relevant sodium contents (x = 0.33, 0.5 and 0.67). The MnNi6 long range ordering observed in Na0.67Mn0.67Ni0.33O2 was preserved in Na0.62Mn0.75Ni0.25O2 as shown through neutron diffraction. These results suggest that the Na+-ion/vacancy ordering is due to

25



charge ordering between consecutive TM planes. We propose that the Mn3+ in Na0.62Mn0.75Ni0.25O2 introduces defects into the Ni-Mn charge order and in turn affects the ordering in the Na-plane. The phase transitions that occur at high (> 4.1 V ; P2-O2) and low voltages (< 2.0 V ; P2P2’) in P2-type materials cause extreme changes to the oxygen lattice that are generally viewed as detrimental to the electrochemical performance. The results presented here on Na0.62Mn0.75Ni0.25O2 ran counter to conventional understanding on the effect phase transitions have on the performance in at least two ways. First, allowing both phase transitions to occur during cycling provided similar capacity retention, and much higher overall capacity than, limiting the cycling voltage window to avoid both transition. When only one of the phase transitions (i.e. either P2-O2 or P2-P2’) was allowed to occur the capacity retention was significantly reduced since the shorter voltage window limited the full recovery of the charge capacity during initial cycling. Second, excellent rate capability was demonstrated even while traversing both phase transitions. The overall performance of Na0.62Mn0.75Ni0.25O2 makes it a viable candidate for use as a cathode material for sodium-ion batteries. Further work is needed to optimize the performance for practical application and to fully understand the improved behavior.

Acknowledgements The submitted manuscript has been created by UChicago Argonne, LLC, Operator of Argonne National Laboratory (“Argonne”). Argonne, a U.S. Department of Energy Office of Science laboratory, is operated under Contract DE-AC02-06CH11357. The U.S. Government retains for itself, and others acting on its behalf, a paid-up, nonexclusive, irrevocable worldwide license in said article to reproduce, prepare derivative works, distribute copies to the public, and

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perform publicly and display publicly, by or on behalf of the Government. The Department of Energy will provide public access to these results of federally sponsored research in accordance with the DOE Public Access Plan. http://energy.gov/downloads/doe-public-access-plan MRCAT operations are supported by the Department of Energy and the MRCAT member institutions. This research used resources of the Advanced Photon Source, a U.S. Department of Energy (DOE) Office of Science User Facility operated for the DOE Office of Science by Argonne National Laboratory under Contract No. DE-AC02-06CH11357. The materials studied herein are based upon work supported by Laboratory Directed Research and Development (LDRD) funding from Argonne National Laboratory, provided by the Director, Office of Science, of the U.S. Department of Energy under Contract No. DE-AC02-06CH11357.

Supporting Information XANES Mn K-edge of Na0.62Mn0.75Ni0.25O2 electrode in different charged states along with three manganese standards and Rietveld refinement of neutron powder diffraction. This information is available free of charge via the Internet at http://pubs.acs.org

References (1) Slater, M. D.; Kim, D.; Lee, E.; Johnson, C. S., Sodium-Ion Batteries. Adv. Funct. Mater. 2013, 23, 947-958. (2) Kim, S.-W.; Seo, D.-H.; Ma, X.; Ceder, G.; Kang, K., Electrode Materials for Rechargeable Sodium-Ion Batteries: Potential Alternatives to Current Lithium-Ion Batteries. Adv. Energy Mater. 2012, 2, 710-721. (3) Yabuuchi, N.; Komaba, S., Recent Research Progress on Iron- and Manganese-Based Positive Electrode Materials for Rechargeable Sodium Batteries. Sci. Technol. Adv. Mater. 2014, 15, 043501. (4) Thackeray, M. M.; Kang, S.-H.; Johnson, C. S.; Vaughey, J. T.; Benedek, R.; Hackney, S. A., Li2MnO3-Stabilized LiMO2 (M = Mn, Ni, Co) Electrodes for Lithium-Ion Batteries. J. Mater. Chem. 2007, 17, 3112-3125. 27



(5) Thackeray, M. M.; Johnson, C. S.; Vaughey, J. T.; Li, N.; Hackney, S. A., Advances in Manganese-Oxide ‘Composite’ Electrodes for Lithium-Ion Batteries. J. Mater. Chem. 2005, 15, 2257-2267. (6) Kim, D.; Kang, S.-H.; Balasubramanian, M.; Johnson, C. S., High-Energy and High-Power Li-Rich Nickel Manganese Oxide Electrode Materials. Electrochem. Commun. 2010, 12, 1618-1621. (7) Johnson, C. S.; Li, N.; Vaughey, J. T.; Hackney, S. A.; Thackeray, M. M., Lithium– Manganese Oxide Electrodes with Layered–Spinel Composite Structures xLi2MnO3·(1−x)Li1+yMn2−yO4 (0

Vacancy Ordering in P2-Type Sodium

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