Effect of Sodium Content on the Reversible Lithium Intercalation into

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Effect of Sodium Content on the Reversible Lithium Intercalation into Sodium-deficient Cobalt-Nickel-Manganese Oxides NaCo Ni Mn O (0.38#x#0.75) with a P3-type of Structure x

1/3

1/3

1/3

2

Svetlana Ivanova, Ekaterina Zhecheva, Rositsa Kukeva, Georgi Todorov Tyuliev, Diana Nihtianova, Lyuben Mihaylov, and Radostina Stoyanova J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.5b11910 • Publication Date (Web): 01 Feb 2016 Downloaded from http://pubs.acs.org on February 2, 2016

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The Journal of Physical Chemistry

Effect of Sodium Content on the Reversible Lithium Intercalation into Sodium-deficient Cobalt-NickelManganese Oxides NaxCo1/3Ni1/3Mn1/3O2 (0.38≤x≤0.75) with a P3-type of Structure Svetlana Ivanova a, Ekaterina Zhecheva a, Rositsa Kukeva a, Georgi Tyulievb, Diana Nihtianova a,c

a

, Lyuben Mihailov d, Radostina Stoyanova a*

Institute of General and Inorganic Chemistry, Bulgarian Academy of Sciences, 1113 Sofia,

Bulgaria. Fax: +359 2 8705024; Tel: +359 2 979 3915; E-mail: [email protected] b

Institute of Catalysis, Bulgarian Academy of Sciences, 1113 Sofia, Bulgaria.

c

Institute of Mineralogy and Crystallography, Bulgarian Academy of Sciences, 1113 Sofia,

Bulgaria d

Faculty of Chemistry and Pharmacy, Sofia University, 1164 Sofia, Bulgaria

ABSTRACT. Layered lithium transition metal oxides with optimized nickel-manganese content are, nowadays, of primary interest as electrode materials for lithium ion batteries, since they are able to deliver a high capacity at a low cost. Herein we report a new class of less expensive cathode materials, which comprise sodium deficient cobalt-nickel-manganese oxides

ACS Paragon Plus Environment

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NaxCo1/3Ni1/3Mn1/3O2 characterized with a layered structure and broad concentration range of sodium solubility. NaxCo1/3Ni1/3Mn1/3O2 oxides are obtained by thermal decomposition of mixed acetate-oxalate precursors, followed by thermal annealing between 700 and 800oC. In the concentration range of 0.33 < x ≤ 0.75, NaxCo1/3Ni1/3Mn1/3O2 oxides assume a layered structure with a three layer stacking (i.e. P3-type of structure). Based on electron paramagnetic resonance spectroscopy operating in the X-band (9.4 GHz), it is found that the charge compensation of Na deficiency is achieved by preferential oxidation of Ni2+ to Ni3+ and Ni4+, while Co and Mn ions retain their oxidation state of 3+ and 4+ within the whole concentration range. The electrochemical performance of NaxCo1/3Ni1/3Mn1/3O2 in model lithium cells is simply controlled by the amount of sodium content in the pristine compositions: a higher reversible capacity is achieved for sodium rich oxides (i.e. 0.75≥x≥0.67), while sodium-poor oxides (i.e. 0.38≤x≤0.50) display a lower reversible capacity and improved cycling stability. The mechanism of the lithium intercalation into NaxCo1/3Ni1/3Mn1/3O2 is discussed on the basis of ex-situ XRD, HRTEM and X-ray photoelectron spectroscopy analyses.

1. Introduction Mixed lithium-transition metal oxides with a layered crystal structure and a nominal composition of Lix(Co/Ni/Mn)yO2 are, nowadays, the most promising cathode materials for lithium ion batteries, since they are able to deliver higher capacity at lower cost in comparison with the conventional LiCoO2 electrode.1,2 Two groups of compounds can be outlined in particular: manganese-rich and nickel-rich oxides.3-7 Manganese-rich oxides belong to the main group of high-energy cathodic materials and they exhibit complex structure, which is not yet well resolved: the structure can be described as an intermixture between layered-monoclinic nanodomains or as a solid solution with presence of a superstructure. In other notation, the general


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layered structure of manganese-rich oxides is stabilized by introducing lithium ions into the transition metal layers yielding a series of Li[Lix(Co/Ni/Mn)1-x]O2 compositions.3 For this reason manganese-rich oxides are also denoted as lithium-rich oxides.4,5 Nickel-rich oxides, LiNi1−xMxO2, form the next group of high-energy cathodic materials.6,7 They display a specific structural feature of intrinsic disordering of Li+ and Ni2+ ions between layers.1,2 The common property of manganese- and nickel-rich oxides is their activation during the first cycle due to a realization of specific structural rearrangement.4-8 This common feature determines several shortcomings of these oxides, associated with their large first irreversible capacity loss, poor thermal stability in delithiated state, cycle life at elevated temperatures and safety problems.4,5,8 These shortcomings have been, to a great extent, overcome by advancing of concentrationgradient layered oxides with Ni-rich interior and Mn-rich surface.6 All these studies show that the best electrochemical properties can be achieved by optimizing the composition of lithiumtransition metal oxides with respect to the nickel and manganese contents. Recently, a new concept for designing low-cost cathode materials has been advanced. This concept is aimed at using directly sodium deficient manganese-based oxides as electrode materials instead of lithium-rich compositions.9-13 There are four reasons for choosing sodium deficient manganese-based oxides. First of all, sodium is more abundant in nature and it has a lower cost in comparison with Li. Second, sodium deficient manganese-based oxides can participate in reactions of intercalation of both Na+ and Li+ ions, as well as in ionic reactions of exchange with Li+ ions.9-15 Similar to lithium-transition metal oxides, sodium-transition metal oxides display relatively good electrochemical performance, when they are used as cathode materials in sodium ion cells.14,15 It is noticeable that lithium and sodium ion batteries operate through the same mechanism comprising the reversible electrochemical intercalation of Li+ and


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Na+.14 Finally, sodium deficient transition metal oxides display a layered structure like that of their lithium analogues.16 Because of the bigger ionic radius of Na+ in comparison with that of Li+, sodium-transition metal oxides possess a variety of structures with respect to the layer stacking and sodium site symmetry.16 In general, sodium-transition metal oxides have crystal structure composed of discrete transition metal oxide layers.16,14 The sodium ions are sandwiched between the transition metal oxide layers in such a way as to occupy octahedral or prismatic sites. Based on the number of the transition metal oxide layers in the unit cell and the sites occupied by Na ions, the structure of sodium-transition metal oxides can be classified as O3-, P3- and P2-types according to the notation, proposed by Delmas et al..16 On the contrary of lithium nickel-rich oxides, the cationic mixture between the layers is not observable in case of sodium nickel-manganese oxides due to the greater ionic radius of Na+.10-14 From electrochemical point of view, the preferable compositions are sodium nickelmanganese oxides. They exhibit a layered structure, which is flexible enabling different layer stacking depending on the Na-to-transition metal ratio. There are two stable structural modifications: O3-type for the stoichiometric sodium oxide NaNi0.5Mn0.5O2 and P2-type for the sodium deficient oxide Na2/3Ni1/3Mn2/3O2.17-21 During the electrochemical desodiation of O3NaNi0.5Mn0.5O2 into Ni0.5Mn0.5O2, there occurs a reversible structural transformation from the hexagonal O3 into the hexagonal P3”-type of structure,17 while a phase transformation from P2 to O2 type of structure proceeds in case of P2-Na2/3Ni1/3Mn2/3O2.20 Under soft chemical conditions, a P3-modification with a composition NaxNi0.5Mn0.5O2, 0.5 ≤ x < 0.75, is isolated.12,13

The

partial

extraction

of

Na+

from

P3-NaxNi0.5Mn0.5O2

resulting

in

Nax~0.3Ni0.5Mn0.5O2 leads to the formation of monoclinically distorted P3-phase, characterized by


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an extremely large interlayer space, while in the completely de-sodiated phase the monoclinic distortion is reduced and the interlayer space is still large.13 The flexibility of the layered structure matrix gives rise to a capability of sodium deficient nickel-manganese oxides to intercalate reversibly both sodium as well as lithium ions, a property that controls their electrochemical performance as cathode materials in sodium and lithium ion cells.11,13 All of the sodium nickel-manganese oxides display a good rate capability when they are used in sodium ion cells.11-13,20,21 The electrochemical properties of P2-Na2/3Ni1/3Mn2/3O2 are close to those of P3-NaxNi1/2Mn1/2O2.13 The P2-Na2/3Ni1/3Mn2/3O2 phase displays higher average working potential and volumetric capacity in comparison with the O3-NaNi0.5Mn0.5O2 phase when they are used in sodium ion cells.21 In the case of lithium ion cells, P3-NaxNi0.5Mn0.5O2 phase is only able to intercalate lithium reversibly, which is in contrast to the P2Na2/3Ni1/3Mn2/3O2 phase.13 The electrochemical intercalation of Li+ into P3-NaxNi0.5Mn0.5O2 leads to a structural transformation from the P3- into the O3-type of structure.13 The in-situ generated O3-phase containing simultaneously lithium and sodium determines the further electrochemical response of P3-NaxNi0.5Mn0.5O2 in terms of voltage profile, cycling stability and rate capability.12,13 This is a unique ability of sodium-deficient nickel-manganese oxides with a P3-type of structure for their application as low-cost electrode materials both in sodium and in lithium ion batteries.11,12 Another advantage of using nickel containing layered oxides in sodium ion batteries is their limited reactivity towards CO2 during exposure to ambient atmosphere.22 Searching for high-energy density, low cost and stable cathode materials, mixed lithium cobalt-nickel-manganese oxide, where Co, Ni and Mn ions are in equal amounts, have attracted recent intensive research efforts.23-26 The oxide has a general composition LiCo1/3Ni1/3Mn1/3O2 and it assumes an O3-type of structure. This oxide possesses a high reversible capacity above 4


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V (i.e. about 200 mAh/g) and it displays a good rate capability in comparison with Co-free analogue. Similar to LiCo1/3Ni1/3Mn1/3O2, the sodium analogue NaCo1/3Ni1/3Mn1/3O2 assumes also O3-type of structure.27 In sodium ion cells, NaCo1/3Ni1/3Mn1/3O2 is able to intercalate reversibly 0.5 mol of sodium between 2.00 and 3.75 V, corresponding to capacity of 120 mAh/g.27 Contrary to LiCo1/3Ni1/3Mn1/3O2, the desodiation process proceeds through series of structural transformations including the sequence O3 O1 P3 P1.27 In contrast to Co-free oxides, sodium deficient cobalt-containing oxides NaxCo1/3Ni1/3Mn1/3O2 are not yet well examined in respect to synthesis, structure characterization and their ability to intercalate alkaline ions. Herein we investigate for the first time, the formation of sodium-deficient oxides NaxCo1/3Ni1/3Mn1/3O2. Then the effects of sodium contents on the structure and intercalation properties of NaxCo1/3Ni1/3Mn1/3O2 are investigated. The studies are focused on the capability of NaxCo1/3Ni1/3Mn1/3O2 to be used as cathodes in lithium ion batteries. We applied a simple precursor-based method for the synthesis of oxides, comprising thermal decomposition of mixed acetate-oxalate precursors. The structure and morphology of NaxCo1/3Ni1/3Mn1/3O2 are determined by powder X-ray diffraction and TEM analysis. The oxidation states of cobalt, nickel and manganese ions are analyzed by electron paramagnetic resonance spectroscopy operating in the X-band (9.4 GHz). The lithium intercalation into NaxCo1/3Ni1/3Mn1/3O2 is carried out in model two-electrode lithium cells of the type Li|LiPF6(EC:DMC)|NaxCo1/3Ni1/3Mn1/3O2. The structural and morphological changes during the lithium intercalation are followed by ex-situ XRD and TEM analysis. The surfaces of pristine oxides and that of cycled electrodes are evaluated by X-ray photoelectron spectroscopy (XPS).


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2. Experimental Synthesis of Oxides. The NaxCo1/3Ni1/3Mn1/3O2 oxides were obtained by oxalate-acetate precursor method. According to this method, sodium hydroxide and oxalic acid were mixed at a molar ratio of 1:1 and ground in an agate mortar until the mixture became sticky. Then solid manganese,

nickel

and

cobalt

acetates

were

added,

the

molar

ratio

being

Na:Co:Ni:Mn=x:1/3:1/3:1/3. The nominal sodium content varied between 0.33 and 1.0. The solid residue was heated at 400oC, followed by thermal annealing at 700 and 800°C for 10 hours.

Structure and Morphology Characterization. The X-ray structural analysis was made on Bruker Advance D8 powder diffractometer with CuKα-radiation. Step-scan recordings for structure refinement by the Rietveld method were carried out using 0.02° 2θ steps of 4-s duration. The diffractometer point zero, the Lorentzian/Gaussian fraction of the pseudo-Voigt peak function, the scale factor, the unit cell parameters, the thermal factors, and the line halfwidth parameters were determined. The computer FullProf Suite Program (1.00) was used in the calculations.28 The TEM investigations were performed on a JEOL 2100 transmission electron microscope and a JEOL 2100 XEDS: Oxford Instruments, X-MAXN 80T CCD Camera ORIUS 1000, 11 Mp, GATAN at accelerating voltage of 200 kV. The specimens were prepared by grinding and dispersing the powders in acetone by ultrasonic treatment for 6 minutes. The suspensions were dripped on standard holey carbon/Cu grids. The analysis was carried out by the Digital Micrograph software.

Spectroscopic Characterization. The EPR spectra were recorded in the form of first derivative of the absorption signal of a Bruker EMXplus EPR spectrometer in the X-band (9.4 GHz) within the temperature range of 120-450 K.


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Surface Analysis. X-ray photoelectron spectroscopy (XPS) was carried out using ESCALAB MkII (VG Scientific) electron spectrometer at a base pressure of 5x10-10 mbar in the analysis chamber (during the measurement 1x10-8 mbar), using MgKα X-ray source (excitation energy hν=1253.6 eV). The instrumental resolution measured as the full width at a half maximum (FWHM) of the Ag3d5/2, photoelectron peak was about 1 eV. The energy scale was corrected with respect to the C1s - peak maximum at 285 eV for electrostatic charging. The fitting of the recorded XPS spectra was performed, using a symmetrical Gaussian-Lorentzian curve fitting after Shirley-type subtraction of the background.

Electrochemical

Characterization.

The

electrochemical

charge-discharge

of

NaxCo1/3Ni1/3Mn1/3O2 was carried out by using two-electrode cells of the type Li|LiPF6 (EC:DMC)|NaxCo1/3Ni1/3Mn1/3O2. The positive electrode, supported onto an aluminum foil, was a mixture containing 80% of the active composition NaxCo1/3Ni1/3Mn1/3O2, 7.5 % KS 6L graphite (TIMCAL), 7.5 % Super C65 (TIMCAL) and 5 % polyvinylidene fluoride (PVDF).

The

electrolyte was a 1M LiPF6 solution in ethylene carbonate and dimethyl carbonate (1:1 by volume) with less than 20 ppm of water. The lithium electrodes consisted of a clean lithium metal disk with diameter of 18 mm. The electrochemical reactions were carried out using an eight-channel Arbin BT2000 system in galvanostatic mode. The charge and discharge rates were expressed as C/h, where h is the number of hours needed to insert one lithium per formula unit at the applied current intensity. The cells were mounted in a dry box under Ar atmosphere. The cell was cycled between 4.4 and 1.8 V at C/100, C/44, C/20, C/10 and C/5 rate. The structural changes in the electrode compositions during reversible intercalation and deintercalation were analyzed with lithium half-cells stopped at selected potentials. The electrochemical cells were disassembled insight a glove-box, followed by removing and washing


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of the working electrodes with EC. The electrode samples were covered with parafilm for the XRD experiments in order to avoid the water contamination. The specimens were dispersed in acetone for the TEM experiments and then the suspensions were dripped on standard holey carbon/Cu grids.

3. Results and Discussions Crystal structure of NaxCo1/3Ni1/3Mn1/3O2 (0.38≤x≤0.75) Figures 1(a-e) show the XRD patterns of sodium-deficient cobalt-nickel-manganese oxides NaxCo1/3Ni1/3Mn1/3O2 for x varying between 0.33 and x=0.75. All diffraction patterns consist of a main layered phase having a P3-type of structure and impurities of NiO-like phase. It is of importance that the amount of the impurity NiO-like phase is less than 1% and it is insensitive to the sodium content. The synthesis of NaxCo1/3Ni1/3Mn1/3O2 bears a resemblance with the formation of high-voltage LiNi1/2Mn3/2O4 spinel, where the impurity NiO-like phase pursues always the target phase.29 In case of Na-poor compositions (i.e. 0.33 < x ≤ 0.38, Figs. 1a, 1b), an additional spinel phase grows in intensity. When the sodium content is higher than 0.75, impurity due to Na2CO3 becomes visible (not shown). The appearance of additional spinel and sodium carbonate phases indicates a stability of P3-NaxCo1/3Ni1/3Mn1/3O2 oxides within a limited concentration range: 0.38 ≤ x ≤ 0.75. It is worth mentioning that, in case of sodium-deficient cobalt-nickel-manganese oxides, the concentration limit of stability of the P3-phase is wider in comparison with that for cobalt-free oxides P3-NaxNi1/2Mn1/2O2 (0.50≤ ≤x≤ ≤0.72).12


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Figure 1. XRD patterns of NaxCo1/3Ni1/3Mn1/3O2 oxides with nominal sodium content x=0.33 (a); x=0.38 (b); x=0.50 (c); x=0.67 (d) and x=0.75 (e). The Bragg’s reflections for layered P3phase (R3m space group), NiO-like and spinel-like phases are indicated below the XRD patterns. Red lines correspond to the simulated XRD patterns using Rietveld refinement. The XRD pattern of the main layered phase is interpreted within the framework of a structural model on the basis of a P3-type of structure. There are two other alternative models using space groups R-3m (No. 166) and R3m (No. 160), respectively. For the high-symmetry R3m model, sodium, transition metal ions and oxygen reside in 6c, 3a and 6c positions (0, 0, zNa), (0, 0, 0) and (0, 0, zO), while for the low-symmetry R3m model sodium and oxygen occupy two types of 3a positions including (0, 0, zNa) / (1/3, 2/3, zNa) and (0, 0, zO1) / (0, 0, zO2), and Co/Ni/Mn ions located in one 3a site (0, 0, 0). Historically, the P3-type of structure is described in R3m space group for sodium deficient cobaltates NaxCoO2.30 The same space group is also


10

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used for analysis of the crystal structure of NaxCo1/3Ni1/3Mn1/3O2 obtained by electrochemical desodiation process.27 Therefore, we calculated the structure of NaxCo1/3Ni1/3Mn1/3O2 within the framework of R3m space group (Table 1). However, the higher-symmetry structural model including space group R-3m is also given for comparison (Table 1). As one can see, it is difficult to discriminate between the two structural models on the basis of powder data statistics only (Table 1). It is of importance that the two structural models give self-consistent structural parameters for NaxCo1/3Ni1/3Mn1/3O2. Table 1 Lattice parameters and RB for layered P3-NaxCo1/3Ni1/3Mn1/3O2 calculated in the framework of space group R3m and R-3m. Samples

R3m

R-3m

a ± 0.0001, c ± 0.0001, RB Å Å

a ± 0.0001, c ± 0.0001, RB Å Å

x=0.33

2.8330

16.8140

4.51

2.8316

16.8054

4.19

x=0.38

2.8314

16.8076

4.81

2.8316

16.8048

3.62

x=0.50

2.8308

16.7759

6.63

2.8309

16.7709

3.06

x=0.67

2.8412

16.7069

6.34

2.8412

16.7058

5.96

x=0.75

2.8413

16.6789

7.67

2.8413

16.6777

6.27

Table

1

shows

a

concentration

dependence

of

the

lattice

parameters

of

NaxCo1/3Ni1/3Mn1/3O2. Upon increasing the nominal sodium content from 0.33 to 0.75, the cparameter corresponding to the interlayer space gradually decreases, while the a-parameter expressing the distance between metal ions inside the layers seems invariable between x=0.33 and x=0.50, followed by a slight increase above x>0.5. The lattice volume shows a tendency to decrease between x=0.33 and x=0.50, while above x>0.5 there appears a slight increase. The observed decrease in the interlayer space can be related to the screening effect of Na+ on the


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electrostatic repulsion between charged transition metal layers Co1/3Ni1/3Mn1/3O2. This phenomenon is usually observed in case of layered lithium-transition metal-oxides, where the interlayer distance is a result from the interplay between the electrostatic repulsion (i.e. oxygenoxygen repulsion) and the steric effect (i.e. lithium-oxygen attraction).31 In the case of oxides possessing a strongly ionic character, the electrostatic repulsion prevails, while the steric effect dominates in more covalent oxides. The observed decrease in the interlayer space upon increasing the sodium content suggests an ionic character of NaxCo1/3Ni1/3Mn1/3O2. In addition, the variation in the a-parameter reflects the changes in the oxidation state of transition metal ions in a synchronized way with varying sodium content.

Oxidation state of Co, Ni and Mn ions in NaxCo1/3Ni1/3Mn1/3O2 (0.38≤x≤0.75) It is generally accepted that the Na content and the requirement for charge neutrality yield the average oxidation state of the transition metal ions in layered NaxMO2 oxides. For the stoichiometric composition Na1Co1/3Ni1/3Mn1/3O2, it is supposed that Co, Ni and Mn ions exhibit the oxidation states of +3, +2 and +4 as in the case of the lithium analogue LiCo1/3Ni1/3Mn1/3O2.27,32-34 Among transition metal ions, Ni2+ and Mn4+ ions are paramagnetic with spins states of S=1 and S=3/2, while Co3+ is a diamagnetic (low-spin configuration d6, S=0).33,34 After decreasing the Na content, the charge compensation can be achieved by oxidation either of Co or Ni ions. This will result in a generation of new paramagnetic and diamagnetic ions such as Co4+ (low-spin configuration d5, S=1/2), Ni3+ (low-spin configuration d7, S=1/2) and Ni4+ (low-spin configuration d6, S=0). In comparison with Co and Ni ions, it has been found out that Mn ions possess a stable oxidation state of +4.32,35 In order to monitor the sodium content induced changes in the oxidation states of Co and Ni ions in NaxCo1/3Ni1/3Mn1/3O2 series, an EPR spectroscopy study was involved. Figures 2(a-d)


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show the EPR spectra of NaxCo1/3Ni1/3Mn1/3O2 oxides. All spectra consist of single Lorentzian line with a g-factor and a line width depending on the sodium content and on the registration temperature. Figures 3a and 3b compare the temperature dependence of the g-factor and the EPR line width for various NaxCo1/3Ni1/3Mn1/3O2 compositions. At 450 K, all compositions have the same g-factor of g=1.99 (Fig. 3a). Upon cooling down from 450 K to 120 K, the g-factor decreases and this decrease is more pronounced for the Na-rich compositions (i.e. x=0.75 and x=0.67). The EPR line width increases with the nominal sodium content; this trend is valid within the whole temperature interval of spectra registration (Fig. 3b). For the Na-rich compositions, there is a strong line narrowing from 450 to 180 K, while after that the EPR line width reaches its minimal value. In contrast to Na-rich compositions, the EPR line width is slightly dependent on the registration temperature for the Na-poor compositions (i.e. x=0.38 and x=0.50).

Figure 2 EPR spectra at 120 K for NaxCo1/3Ni1/3Mn1/3O2 with nominal sodium content x=0.38 (a), x=0.50 (b), x=0.67 (c) and x=0.75 (d). The black and red lines correspond to the experimental and simulated spectra.


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Figure 3 The temperature dependence of the g-factor (a) and the EPR line width, ∆Hpp (b), for NaxCo1/3Ni1/3Mn1/3O2 with nominal sodium content x=0.38, x=0.50, x=0.67 and x=0.75. The EPR references P2-Na2/3MnO2+δ, O3-LiCo1-2xNixMnxO2 (LCNM) and O3-LiNi1/2Mn1/2O2 (LNM) are also shown. To facilitate the analysis of EPR parameters for NaxCo1/3Ni1/3Mn1/3O2 we used two kinds of EPR references: layered sodium manganese oxides Na2/3MnO2+δ with a P2-type of structure as an EPR reference for exchange coupled Mn4+,36 and Ni,Mn-cosubstituted lithium cobalt oxide LiCo1-2xNixMnxO2 (x=1/3 and ½) with an O3-type of structure as an EPR reference for simultaneously stabilized Mn4+ and Ni2+ ions.25,37 Although all EPR references exhibit EPR signal with a Lonrentzian line shape, the differences between them lie in the values of g-factors and EPR line widths. For the unsubstituted manganese oxide Na2/3MnO2+δ, the exchange coupled Mn4+ ions inside the MnO2-layers give rise to a single Lorentzian line with a g-factor of 1.996(1), which remains constant within the temperature interval from 100 to 450 K (Fig. 3a). The comparison shows that the g-values of exchanged coupled Mn4+ in P2- Na2/3MnO2+δ fit well with the high-temperature values of the g-factor for all NaxCo1/3Ni1/3Mn1/3O2 compositions. This reveals that Mn4+ ions give rise to the EPR signal from NaxCo1/3Ni1/3Mn1/3O2 oxides. For layered O3-LiCo1/3Ni1/3Mn1/3O2 oxides, where Ni2+ and Mn4+ occur simultaneously, an EPR response from Mn4+ ions has only been detected.25 However, the presence of Ni2+ ions has an impact on


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the temperature dependence of g-factor: there is a decrease in the g-factor with a lowering of the registration temperature due to the development of exchange interactions between Mn4+ and Ni2+ ions (Fig. 3a). The temperature-induced changes in the g-factor are more pronounced for the Cofree composition O3-LiNi1/2Mn1/2O2, where a strong anti-ferromagnetic interactions between Ni2+ and Mn4+ ions in transition metal layers are occurring.37 The EPR spectra of O3-Cocontaining LiCo1/3Ni1/3Mn1/3O2 and cobalt-free O3-LiNi1/2Mn1/2O2 oxides reveal the role of diamagnetic ions on the EPR parameters of layered oxides: the greater the number of diamagnetic Co3+ ions appear around paramagnetic Ni2+ and Mn4+ ions, the weaker is the temperature dependence of the g-factor (Fig. 3a). It should be mentioned that the higher values of the g-factor of the lithium analogue LiCo1/3Ni1/3Mn1/3O2 reflect presence of Ni2+ ions in the LiO2-layers, which yield 180o-magnetic interactions between the layers in addition to 90omagnetic interactions inside the layers.25 Taking into account the above role of diamagnetic ions, we can divide Na-poor and Na-rich compositions into two groups, which can be differentiated on the basis of amounts of diamagnetic ions. It appears that Na-poor compositions (x=0.38 to x=0.50) contain more diamagnetic ions in regard to the Na-rich compositions ((x=0.67 to x=0.75). The EPR line width is another EPR parameter providing information on the oxidation state of transition metal ions. The EPR line width is smaller for the unsubstituted manganese oxide Na2/3MnO2+δ, where exchange interactions between similar Mn4+ ions are dominant (Fig. 3b). When Ni2+ ions occur in close proximity to Mn4+, the EPR signal becomes broader (Fig. 3b). The main difference between Mn4+oxide and Ni2+-Mn4+-containing oxides is the temperature dependence of the EPR line width (Fig. 3b). The EPR line width of Mn4+ ions is slightly temperature dependent in case of sodium manganese oxide P2-Na2/3MnO2+δ, while the exchange


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interactions between dissimilar paramagnetic Ni2+ and Mn4+ ions contribute to a strong temperature variation of the EPR line width for the co-substituted Ni-Mn oxides like the system LiCo1-2xNixMnxO2. Furthermore, the appearance of diamagnetic Co3+ ions yields a tuning in the temperature dependence of the EPR line width due to the dilution of the magnetic Ni2+-Mn4+ spin system. Based on these correlations, two specific EPR features can be distinguished for Na-rich and Na-poor compositions NaxCo1/3Ni1/3Mn1/3O2. First, the broader signals and the stronger temperature dependence of the line width are an indication of appearance of larger number of paramagnetic ions around Mn4+ for the Na-rich compositions (x=0.75 and x=0.67). This can be achieved by preferential oxidation of paramagnetic Ni2+ to paramagnetic Ni3+ instead of oxidation of diamagnetic Co3+ to paramagnetic Co4+, since the temperature induced changes in the g-factor values bears a resemblance with that for the P2-Na2/3Ni1/3Mn2/3O2, where only Ni2+ and Mn4+ are appearing. The second EPR feature is related to the oxidation states of transition metal ions for Na-poor compositions (x=0.38 and x=0.50): both the narrower signals and the weaker temperature dependence of the EPR line width indicate a lack of dissimilar paramagnetic ions around the target Mn4+ ions. This can be realized by further oxidation of Ni2+ and Ni3+ into diamagnetic Ni4+ together with preserving diamagnetic Co3+ ions. In summary, the charge compensation in sodium deficient oxides NaxCo1/3Ni1/3Mn1/3O2 is achieved by the preferential oxidation of Ni2+ into Ni3+ and Ni4+ ions, while Co and Mn ions retain their stable oxidation states of +3 and +4, respectively.

Reversible Li intercalation into P3-NaxCo1/3Ni1/3Mn1/3O2 (0.38 ≤ x ≤ 0.75) The sodium content induced changes in the oxidation state of transition metal ions are clearly demonstrated by the open circuit voltage (OCV) of NaxCo1/3Ni1/3Mn1/3O2. Upon increasing the sodium content from 0.38 to 0.75, there is a gradual decrease in the OCV (Table 2). This is an


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electrochemical evidence for increasing the average oxidation state of transition metal ions by lowering the sodium content. Table 2 Open circuit voltage (OCV) for NaxCo1/3Ni1/3Mn1/3O2, as well as first discharge capacity, amount of intercalated Li+ and mean oxidation state (OS) of transition metal ions calculated at C/100 and C/20 rates. Nominal OCV, Discharge Capacity, Compositions V mAh/g (C/100)

Amount of OS intercalated (C/100) Li+ (C/100)

Discharge Capacity, mAh/g (C/20)

Amount of intercalated Li+ (C/20)

x=0.38

3.901

140

0.51

3.51

123

0.45

x=0.50

3.887

132

0.49

3.49

102

0.38

x=0.67

3.799

90

0.35

3.35

73

0.28

x=0.75

3.706

95

0.38

3.38

65

0.20

Figures 4a and 4b compare the first discharge curves for NaxCo1/3Ni1/3Mn1/3O2 registered at two rates: C/20 and C/100, respectively. In order to outline the voltage profiles, where the electrochemical reaction takes place, the same figure gives also the first derivative of the discharge curve. At low and high rates, all compositions deliver a capacity that depends on the sodium content. As it can be expected, the discharge capacity is higher when the discharge rate is lower. Assuming lack of side redox reactions, the first discharge capacity determined at a C/100 rate can serve as a measure for the average oxidation state of transition metal ions. Table 1 shows the calculated amount of intercalated Li+ per formula unit NaxCo1/3Ni1/3Mn1/3O2, as well as the corresponding oxidation states of the transition metal ions. The lower the sodium content is, the higher is the amount of intercalated Li+. With the exception of the oxide having x=0.75, the amount of intercalated Li+ compensates for the sodium deficiency in NaxCo1/3Ni1/3Mn1/3O2. This is consistent with OCV of NaxCo1/3Ni1/3Mn1/3O2 (Table 2). For the oxide with content x=0.75, the amount of intercalated Li+ exceeds the sodium deficiency, which implies occurring of some


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interaction of the electrode composition with Li electrolyte solution. At a high rate of discharge, the amount of intercalated Li is always lower than that expected for the formula unit (Table 2).

Figure 4. First discharge (a,b) and charge (c,d) curves and their first derivatives for NaxCo1/3Ni1/3Mn1/3O2 registered at C/20 (a,c) and C/100 (b,d) rates. The next issue is related to the shape of the discharge curve (Figs. 4a and 4b). For all the oxides NaxCo1/3Ni1/3Mn1/3O2, the discharge curve displays two peaks: a peak between 4.0 and 3.5 V and another one between of 3.3 and 2.3 V, respectively (Figs. 4a and 4b). At low discharge rate (i.e. C/100), the Na-rich compositions display two well resolved peaks at 3.7 and 2.9 V. Contrary to Na-rich compositions, the discharge curve profiles are more complicated for the Napoor compositions: the high-voltage peak is broader with a center of gravity at about 3.8 V, and the low-voltage peak is split into two components at about 2.8 and 2.6 V. At high discharge rate (i.e. C/20), the high-voltage peak is less resolved especially for the Na-rich composition, while


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the low-voltage peak remains with a slight shift in the centre of gravity. It is of importance that the capacity associated with the high-voltage peak is lower in comparison with that of the lowvoltage peak irrespective of the discharge rate used: at a rate of C/100, the oxide with x=0.38 intercalates 0.07 and 0.44 mol of Li+ at 3.8 V and below 2.8 V, while at a rate of C/20 these amounts are 0.12 and 0.26, respectively. The observed behavior is obeyed for the oxides with higher sodium content (Figs. 4a and 4b). The variations in the shape of the first discharge curve are related with the sodium content induced changes in the oxidation state of transition metal ions in NaxCo1/3Ni1/3Mn1/3O2. Based on the oxidation state of transition metal ions, determined by EPR analysis, one can expect that Li+ intercalation into NaxCo1/3Ni1/3Mn1/3O2 would take place at the expense of the reduction of Ni and Mn ions only. The reduction of Co3+ into Co2+ is usually energetically less favorable process for the layered transition lithium metal oxides.38 To check the above supposition, we take into account the electrochemical process of lithium intercalation into Cofree NaxNi0.5Mn0.5O2 analogue.12,13 In the case of x=0.50 composition, two well separated voltage plateaus at 3.72 and 3.05 V have been observed and they have been attributed to a consecutive reduction of nickel and manganese ions during lithium intercalation.12,13 The comparison shows that the high-voltage peak is one and the same for Co-free NaxNi0.5Mn0.5O2 and Co-containing NaxCo1/3Ni1/3Mn1/3O2 oxides, which allows us to assign the plateau at 3.72 V to reduction of highly oxidized Ni3+/4+ ions into Ni2+. The low-voltage plateau at 2.70 V is due to reduction of Mn4+ into Mn3+. The observed shift in the low-voltage plateau for Co-containing NaxCo1/3Ni1/3Mn1/3O2 oxides (i.e. from 3.05 to 2.70 V) implies that Co3+ ions participate also in the electrochemical reaction in addition to Mn4+ ions. This is an unexpected result considering the electrochemical properties of layered oxides based on lithium, cobalt, nickel and manganese:


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the Co3+/Co2+ redox couple becomes usually active below 2.0 V and this process is accompanied by a structural transformation.39,40 Contrary to lithium-containing oxides, it has recently been found out that sodium-containing oxides of Na2/3Co2/3Mn1/3O2 composition are able to intercalate Na+ at the expense of the simultaneous reduction of Co3+ and Mn4+ ions into Co2+ and Mn3+.41 This finding permits us to relate the low-voltage peak of NaxCo1/3Ni1/3Mn1/3O2 to reduction of Mn4+ into Mn3+ together with reduction of Co3+ into Co2+. In addition, the reduction of Co ions proceeds more easily in case of Na-poor compositions (i.e. 0.38 ≤ x ≤ 0.50), leading to a complex discharge curve profiles. This is an interesting phenomenon, which deserves further experimental and theoretical investigations. After the reverse process, the charge curve becomes more complicated in comparison with the discharge one (Figs. 4c and 4d). At least three peaks can be distinguished. The plateau at 3.1 V is well defined especially with the Na-poor compositions. The capacity corresponding to this peak is smaller in comparison with the one due to the reduction peaks below 2.8 V. Above 3.5 V, there are two broad peaks at about 3.9 and 4.2 V, where most of the 60% of the charge capacity is delivered. The complex shape of the charge curves implies that all Ni2+/Ni3+/4+, Co2+/Co3+, Co3+/Co4+ and Mn3+/Mn4+ redox pairs participate in the electrochemical reaction. In addition, the occurrence of side reactions (such as electrode-electrolyte interactions, electrolyte decomposition, etc.) at potentials higher than 4.2 V cannot be excluded, especially in the case of Na-rich compositions. It is of importance, that all the oxides deliver a charge capacity corresponding to an extraction of more than 0.7 mol alkaline ions (Figs. 4c and 4d). This amount exceeds significantly the content of inserted lithium ions, thus giving some evidence for a simultaneous extraction of Na+ and Li+ from oxides during the first charge. The extraction of both types of alkaline ions is also consistent with the complex shape of the charge curves. The


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close inspection of charge curves shows that the charge profile displays a tendency to become one and the same for oxides having different nominal sodium content. A specific feature of the oxides after the first cycle is high irreversibility. The lower the rate of charge and discharge is, the higher is the irreversible capacity (Figs. 5a, 5b and 5c). Furthermore, the irreversibility depends on the nominal sodium content. In case of Na-rich oxides the Coulombic efficiency varies between 40 and 60 % depending on the charge/discharge rate, while the Na-poor compositions display better Coulombic efficiencies with values varying between 80-90% (Fig. 5a, 5b and 5c). Based on the first irreversible capacity, it appears that the Na-poor compositions are better performing, irrespective of the charge/discharge rate used.

Figure 5. Cycling stability of NaxCo1/3Ni1/3Mn1/3O2 having x=0.38, 0.67 and 0.75 using rates of C/10 (a), C/20 (b) and C/44 (c). The blank and full symbols indicate the charge and discharge capacities. The inset gives the corresponding Coulombic efficiency. Further cycling causes a drastic decrease in the first irreversible capacity of all oxide (Fig. 5a, 5b and 5c). A stable electrochemical performance is achieved after 2 to 5 cycles (Fig. 5a, 5b and 5c). Figures 6(a-f) compare the charge/discharge curves of NaxCo1/3Ni1/3Mn1/3O2 after 2, 5 and 10 cycles at rates of C/44, C/20 and C/10. Two features can be outlined: (i) the Coulombic efficiency reaches 99-100% after 5 cycles with all the oxides, and (ii) the charge/discharge curves become smoother and insensitive in regard to the nominal sodium content. This fact supposes a structural transformation of NaxCo1/3Ni1/3Mn1/3O2 oxides during the electrochemical


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charge/discharge processes, as a result of which the peculiarities of oxides having different nominal sodium contents are smeared.

Figure 6. Charge/discharge curves after 2, 5 and 10 cycles at a rate of C/44 (a,d), C/20 (b,e) and C/10 (c,f) for NaxCo1/3Ni1/3Mn1/3O2 with x=0.38 (a,b,c) and 0.75 (d,e,f). However, the differences inherited from the pristine compositions are manifested in the reversible capacity. The cycling stabilities of Na-poor and Na-rich oxides are compared in Figures 5a, 5b and 5c. The reversible capacity shows a dependence on the nominal sodium content. This dependence is obeyed irrespective of the rate of charge/discharge used. The reversible capacity is higher in case of oxide with x=0.75, while the oxide with x=0.38 displays a lower reversible capacity. In this respect, it appears that the Na-rich compositions are more flexible to insert lithium reversibly in comparison to the Na-poor ones. In addition, the occurrence of impurity spinel phase in the oxide with x=0.38 can also contribute to a lower reversible capacity.


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Structural changes of P3-NaxCo1/3Ni1/3Mn1/3O2 during charge/discharge In order to rationalize the structural changes of P3-NaxCo1/3Ni1/3Mn1/3O2 during the electrochemical process of charge/discharge, we applied ex-situ XRD and TEM techniques. Figure 7 gives the XRD patterns of electrode compositions Na0.67Co1/3Ni1/3Mn1/3O2 obtained after the first discharge down to 2.5 V and further down to 1.8 V, as well as of the oxide Na0.67Co1/3Ni1/3Mn1/3O2 obtained after 8 cycles between 1.8 V and 4.4 V (the cell test is stopped at 1.8 V). The lithium intercalation causes a strong broadening in all diffraction peaks, as a result of which we study the changes in the basal plane only. The diffraction peak due to the basal plane of pristine Na0.67Co1/3Ni1/3Mn1/3O2 disappears at the beginning of lithium intercalation, which indicates complete involvement of the layered oxide in the electrochemical reaction. The first intercalation of Li+ into Na0.67Co1/3Ni1/3Mn1/3O2 leads to a discrete shifting of the diffraction peak: the interlayer space decreases from 5.53 Å to 5.15 Å after the high-voltage peak, followed by a further contraction to 4.92 Å during the low-voltage peak. The observed strong decrease in the interlayer space is related to the intercalation of Li+ having smaller ionic radius at the expense of Na+ with larger ionic radius. The important finding here is the discrete decrease in the interlayer space during Li+ intercalation into P3-NaxCo1/3Ni1/3Mn1/3O2. This indicates structural transformation comprising a change in the layered sequence and alkaline site occupancy in a discrete way. Although Li+ ions prefer to reside octahedral positions in the layered structure, both octahedral and prismatic sites are occupied by Na+ ions. This phenomenon is well demonstrated in the case of Co-free analogue NaxNi1/2Mn1/2O2: it has been found that there is a direct transformation of the P3- into O3-type of structure during the Li+ intercalation.13 After several cycles between 1.8 V and 4.4 V, the interplanar distance of NaxCo1/3Ni1/3Mn1/3O2 tends to reach 4.75 Å, which is lower value in comparison with that for the initially lithiated oxide (i.e.


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4.92 Å). This indicates that P3-NaxCo1/3Ni1/3Mn1/3O2 is transformed into stable lithiated oxide after few cycles.

Figure 7. ex-situ XRD patterns of Na0.67Co1/3Ni1/3Mn1/3O2 electrode after the first discharge to 2.5 V, after first discharge to 1.8 V and after 8 cycles between 1.8-4.4 V (the cell is stopped at 1.8 V). For the sake of comparison, the XRD pattern of the pristine oxide is also shown. The stability of the layered structure during cycling is proved by ex-situ TEM analysis. Figures 8(a-f’) compare the bright field images, SAED and HRTEM of pristine P3Na0.67Co1/3Ni1/3Mn1/3O2 and cycled compositions derived on the basis of Na-poor and Na-rich oxides (i.e. x=0.38 and x=0.67) after 8 cycles between 1.8 V and 4.4 V (the cell test is stopped at 1.8 V, i.e. the oxide is in a completely lithiated state). Pristine Na0.67Co1/3Ni1/3Mn1/3O2 consists of well-faceted thin particles with dimensions varying between 20 and 100 nm (Fig. 8a). The thin particles with the same sizes remain intact after the electrochemical reaction carried out with both Na-rich and Na-poor oxides (Figs. 8b and 8c). The polycrystalline electron diffraction pattern of pristine and cycled oxides is indexed on the basis of phase mixture between layeredand NiO-like phases. The difference between pristine and cycled oxides originates from the lattice parameters

of

the layered

phase:

the lattice parameters

for pristine P3-

Na0.67Co1/3Ni1/3Mn1/3O2 are a=2.837 Å and c=16.735 Å and, after the electrochemical charge/discharge, they become a=2.84 Å and c=14.44 Å. The observed reduction in the c-


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parameter reveals a strong interlayer contraction (i.e. from 5.7 to 4.8 Å), while the small variation of the a-parameter is associated with a slight alteration of the distance between metal ions inside the layers. Contrary to the layered phase, the NiO-like phase has unchanged lattice parameters. This proves that only the layered phase participates in the electrochemical reaction, while the NiO-like phase is electrochemically inert.

a

b

c


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Figure 8. Bright field micrographs, polycrystalline electron diffraction patterns and SAED for pristine Na0.67Co1/3Ni1/3Mn1/3O2 (a) and NaxCo1/3Ni1/3Mn1/3O2 electrodes with nominal x=0.67 and x=0.38 (b and c) cycled 20 times between 1.8 and 4.4 V at a C/20 rate. HRTEM of pristine Na0.67Co1/3Ni1/3Mn1/3O2 (d) and corresponding cycled electrodes NaxCo1/3Ni1/3Mn1/3O2 with x=0.67 (e) and x=0.38 (f and f’). The strong decrease in the c-parameter is also observed by SAEDs of thin particles: for the pristine oxide, the lattice parameters estimated by SAED, taken along the [0 1 0] direction, are a=2.84 Å and c=16.74 Å, while in case of cycled compositions derived from Na-rich and Napoor oxides, SAED taken along [-1 1 1] directions gives the same lattice parameters of a=2.84 Å and c=14.4 Å (Figs. 8a, 8b and 8c). It is worth mentioning that lattice parameters, evaluated by SAED, coincide with those determined from the polycrystalline patterns (Table 1). The strong deviation of the lattice parameters of cycled compositions from those of pristine oxides is related to formation of layered lithium-containing phase during reversible lithium intercalation. For the sake of comparison, the lithium analogue LiCo1/3Ni1/3Mn1/3O2 (obtained by a solid state reaction at 950 oC) having an O3-type of structure exhibits lattice parameters of a=2.8541 Å and c=14.2130 Å.25 On the one hand, the comparison indicates transformation from P3- into O3-type of structure during cell cycling. On the other hand, the higher value of the c-parameter for the sodium-derived oxides suggests an in-situ formation of layered phases containing simultaneously both lithium and sodium ions. Using BF-STEM, the composition of the layered phase P3-NaxCo1/3Ni1/3Mn1/3O2 is evaluated after the electrochemical reaction (Fig. 9). The composition of the pristine oxide with


26

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nominal sodium content x=0.67 is Na0.56Co0.34Ni0.31Mn0.34O2 and after the electrochemical reaction the sodium content decreases from 0.56 down to 0.03 mol, while Co, Ni and Mn ion contents remain unchanged: 0.34 mol for Co, 0.30 mol for Ni and 0.35 mol for Mn (Fig. 9a). In addition, the Co, Ni, Mn and O elements are homogeneously distributed over nano-sized particles, subjected to the electrochemical charge/discharge (Fig. 9a). The constancy in the ratio between Co, Ni and Mn ions, as well as their homogeneous distribution over particles, give clear evidence for the stability of transition metal layers during lithium intercalation process. It is of importance that the composition of the lithiated oxide, derived from the Na-rich oxide, is close to that derived from the Na-poor oxide: 0.03 mol for Na, 0.35 mol for Co, 0.35 mol for Ni and 0.30 mol for Mn (Fig. 9b). This fact proves that, irrespective of the nominal sodium content, all the oxides undergo changes in their chemical compositions, culminating after few cycles in formation of the lithiated oxides with the same composition. Taking into account the structural requirement for restricted site occupancy by alkaline ions, we can suppose that in-situ formed oxides exhibit such a composition, where both lithium and sodium ions coexist in alkaline layers: Liy1Nax~0.03Co0.35Ni0.35Mn0.30O2. The same picture has been observed during the reversible intercalation of lithium into Co-free analogue NaxNi1/2Mn1/2O2.13

(a)


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(b)

Figure 9. BF-STEM images and corresponding composition map of CoKα1, NiKα1, MnKα1, O Kα1 and Na Kα1-2 for cycled electrodes NaxCo1/3Ni1/3Mn1/3O2 with nominal sodium content of x=0.67 (a) and 0.38 (b). All the results lead to the conclusion that the electrochemical reaction of intercalation of Li into NaxCo1/3Ni1/3Mn1/3O2 takes place by in-situ formation of layered lithium-containing phase accompanied by a transformation from the P3-type into the O3-type of structure. Lithiated oxides derived from Na-poor and Na-rich compositions display well crystallized particles


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without any preferred orientations. Although both cycled systems display similar chemical compositions, the difference between them lies in the structural homogeneity of the in-situ formed O3-phase (Figs 8e, 8f and 8f’). HRTEM images of selected particles show that the interlayer space varies between 4.7 and 4.8 Å for the lithiated oxide, obtained from sodium-poor oxide (i.e. x=0.38, Figs. 8f and 8f’), while the sodium-rich oxide yields lithiated oxide with a constant interlayer space of 4.8 Å (within the range of experimental error, Fig. 8e). For the sake of comparison, the pristine P3-phase displays individual nano-particles with a constant interlayer space of 0.56 nm (Fig. 8d). These d-values are in good agreement with the lattice parameters, determined by SAED and polycrystalline diffraction patterns (Figs. 8a, 8b and 8c). The formation of domains with variable interlayer space can be related to the improved cycling stability of compositions derived from Na-poor oxides, while compositions derived from the Narich oxides are characterized by a homogeneous domain structure and a higher reversible capacity.

Surface composition of NaxCo1/3Mn1/3Ni1/3O2 electrodes The next important issue is related to the surface of oxides, which are in contact with LiPF6based electrolyte. The thermal instability of LiPF6 determines its high reactivity with respect to surface of lithium transition metal oxides.42-44 Even at room temperature, LiPF6 has been shown to decompose into LiF and PF5, the latter one being a strong Lewis acid.42-44 The reaction product PF5 is mainly responsible for the formation of the solid electrolyte interphase layer that is composed of mixed organic and inorganic compounds on the oxide surface.42-44 In order to monitor the changes on the surface of NaxCo1/3Mn1/3Ni1/3O2, Figures 10(a-d) show the XPS spectra of electrode compositions in the energy region characteristic of sodium, lithium, fluorine and phosphorus. Three types of electrode compositions were studied:


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(a)

(c)

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(b)

(d)

Figure 10 XPS spectra in the Na1s (a), Na 2s and Li 1s (b), F 1s (c) and P 2p (d) regions for pristine Na0.67Co1/3Ni1/3Mn1/3O2 and electrode Na0.67Co1/3Ni1/3Mn1/3O2 after the first discharge to 2.5 V, after the first discharge to 1.8 V and after 8 cycles between 1.8 and 4.4 V (the cell is stopped at 1.8 V). Na0.67Co1/3Mn1/3Ni1/3O2 after the first discharge down to 2.5 V, Na0.67Co1/3Mn1/3Ni1/3O2 after the ACS Paragon Plus Environment

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first discharge down to 1.8 V, and Na0.67Co1/3Mn1/3Ni1/3O2 after the 8 cycles between 1.8 and 4.4 V and stopped at 1.8 V. We used both Na1s and Na2s spectra for the surface analysis in order to probe different sampling depth (Figs. 10a and 10b): Na1s provides information about the few top surface layers only, while Na2s signal can be regarded as more “bulk-phase-like”information. In addition, the electron energy for Na2s falls within the range, where the electron energy for Li1s appears (Fig. 10b). This means that the surface sensitivity to Na2s and Li1s is similar. In the energy region of Na1s and Na2s, all the cycled electrode compositions display broad signals with binding energies varying between 1072.3 and 1072.6 eV, as well as between 64.1 and 64.5 eV, respectively (Figs. 10a and 10b). These values are shifted with more than 1 eV in comparison with those for the pristine powder Na0.67Co1/3Mn1/3Ni1/3O2 (i.e. 1070.6 and 62.6 eV, respectively, not shown). The binding energies for both Na1s and Na2s peaks fall within the range, where the binding energies for Na in NaF occur at 1072 eV for Na1s and 63 eV for Na2s, respectively.45 This allows assigning the two Na1s and Na2s peaks to NaF. The spreading of NaF all over the surface of Na0.67Co1/3Mn1/3Ni1/3O2 electrode starts during the electrode fabrication. This can be explained in terms of interaction of oxide surface with PVDF used as a binder. It is interesting to note that the high reactivity towards PVDF has also been found out for Co-free analogue Na0.67Mn1/2Ni1/2O2, while polyanion compounds such as sodium manganese phospho-olivines are inert.12 The surface-formed NaF remains stable during the cell cycling. When Na0.67Co1/3Mn1/3Ni1/3O2 is being discharged down to 2.5 V, a signal due to Li1s becomes visible (Fig. 10b). The binding energy of Li1s amounts to 56.1 eV, which is typical of LiF. For the sake of comparison, the binding energy of Li1s in the lithium analogue LiCo1/3Ni1/3Mn1/3O2 is 54.0 eV, i.e. with 2 eV lower than that of LiF (not shown). The surface ratio of Li-to-Na calculated based on Li1s and Na2s is 0.2, which indicates that LiF is deposited


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immediately after the first discharge down to 2.5 V. Further discharge to 1.8 V causes only an increase in the Li-to-Na ratio from 0.2 to 1.8, while the binding energy of Li1s seems to remain unchaged. During the cell cycling, the XPS spectrum displays the Li1s peak with the same binding energy, while the Na2s peak is not distinguished. However, the Na1s peak is still well visible. This indicates an enrichment of the oxide surface in LiF, which remains stable during the cell cycling (Fig. 10b). On the other hand, the appearance of the Li1s signal gives a clear evidence for interaction of Li+ with Na0.67Co1/3Mn1/3Ni1/3O2. The F1s and P2p spectra display, furthermore, the changes occurring on the oxide surface (Figs. 10c and 10d). The spectrum of pristine electrode in the region of F1s consists of two overlapping signals: a strong signal at 688.2 eV and a less intensive signal at 685.1 eV (Fig. 10c). The relative intensities of the high and low energy signals are 0.82 to 0.18, respectively. The high-energy peak (i.e. at 688.1 eV) is associated with fluorine atoms in the PVDF binder,46 while the low-energy peak (i.e. 685.2 eV) corresponds to fluoride ion in NaF.47 This confirms the observation of NaF on the surface of Na0.67Co1/3Mn1/3Ni1/3O2 electrode during the electrode fabrication. As it is in the case of the pristine electrode, two overlapping signals give rise to the profile of the F 1s spectrum for all discharged and cycled electrodes. The binding energy of the high-energy peak slightly decreases from 688.1 to 687.6 eV, when passing over from the pristine to the cycled electrodes, while the low-energy peak increases by the same order of magnitude: from 685.1 to 685.6 eV. The F 1s peak with a binding energy centred in the range of 687-688 eV can be attributed to fluorine-based compounds such as LixPFyOz and LixPFy,48,49 while the F1s peak is varying between 685 and 686 eV and it can be related to fluoride-based compounds such as LiF and NaF.50 The relative intensity of the two signals is undergoing a strong change during the electrochemical reaction (Fig. 10c): during the cell discharge, there is an increase in the


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relative intensity of the low energy signal due to fluoride-based compounds at the expense of the high energy signal due to the fluorine-based compounds, reaching a value of 0.68 at oxides discharged to 1.8 V. The observed changes in fluoride-fluorine surface compositions result from a dynamic process and after few cycles a stable coverage is obtained: the relative intensities of the high and low energy signals are 0.57 and 0.43, respectively. The P2p spectra of all the cycled electrode compositions consist of one broad peak with a centre of gravity located between 134.5 and 134.7 eV (Fig. 10d). According to the literature data, the peak having the lower binding energy (i.e. at 134.5 eV – with LixPFy.48,49,51,52 Taking into account all of these data, one can identify LixPFy and/or NaxPFy compounds on the surface of the cycled electrode compositions. It appears that mixed alkaline phospho-fluorine compounds are formed immediately after the start of cell discharge and they remain stable during cycling. The appearance of surface LixPFy and/or NaxPFy compounds are also in accordance with a slight variation in the binding energies of both Li1s and Na2s for Na0.67Co1/3Mn1/3Ni1/3O2 electrodes. The surface P-to-F ratio is about 3.0 for all cycled compositions. The formation of surface layers containing Li, Na, P and F prevents the clear observation of XPS spectra due to transition metal ions Mn, Co and Ni. This is demonstrated in Figures 11(a-c), where only the XPS spectra of cycled compositions in the spectra range of Mn 2p3/2 and Mn 2p1/2, Co 2p3/2 and Co 2p1/2 and Ni 2p3/2 are presented. For the sake of comparison, the XPS references for Mn4+, Co3+ and Ni2+ ions in O3-LiCo1/3Ni1/3Mn1/3O2 are also shown. The comparison shows that Mn4+ ions retain their oxidation state (Fig. 11c), while Co ions occur in both +2 and +3 oxidation states (Fig. 11a). In the case of Ni ions, the binding energy at 856.1 eV together with a satellite at 862.4 eV cannot be attributed simply to Ni3+ ions (Fig. 11b). It should


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be taken into account, that Ni2+ ions in NiF2 are causing an appearance of the Ni 2p3/2 signal with a binding energy of about 856-857 eV.53

a

b

c

Figure 11 XPS spectra in the Co (a), Ni (b) and Mn (c) regions for electrode Na0.67Co1/3Ni1/3Mn1/3O2 after 8 cycles between 1.8 and 4.4 V (the cell is stopped at 1.8 V). For the sake of comparison, the XPS references for Co3+, Ni2+ and Mn4+ ions in O3LiCo1/3Ni1/3Mn1/3O2 are also given

4. Conclusions Sodium deficient cobalt-nickel-manganese oxides are obtained by thermal decomposition of mixed acetate-oxalate precursors, followed by thermal annealing between 700 and 800oC. In the concentration range of 0.33 < x ≤ 0.75, NaxCo1/3Ni1/3Mn1/3O2 oxides are forming with a P3-type of structure. By increasing sodium content, the interlayer space gradually decreases, while the distance between metal ions inside the layers seems unchanged. The charge compensation of Na deficiency is achieved by preferential oxidation of Ni2+ to Ni3+ and Ni4+, while Co and Mn ions retain their oxidation state of 3+ and 4+ within the whole concentration range. The sodium-induced increase in the oxidation state of transition metal ions leads to a corresponding enhancement of the OCV of NaxCo1/3Ni1/3Mn1/3O2. During the first discharge at low rates, NaxCo1/3Ni1/3Mn1/3O2 oxides manifest capacities that correspond to the intercalation of lithium in amounts compensating the sodium deficiency. The lithium insertion into


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NaxCo1/3Ni1/3Mn1/3O2 is accomplished at two potential plateaus owing to a selective reduction of transition metal ions: at 3.7 V, there is a reduction of highly oxidized Ni3+/4+ into Ni2+, while below 2.8 V occurs reduction of Mn4+ into Mn3+. The reduction of manganese ions is concomitant with a partial reduction of Co3+ into Co2+, a phenomenon which is generally considered to be rare for layered oxides. The step-wise insertion of Li ions into NaxCo1/3Ni1/3Mn1/3O2 is also manifested by a structural transformation consisting of discrete decrease in the interlayer space from 5.53 to 5.15 Å and from 5.15 to 4.92 Å, respectively. After the reverse process of charging, both lithium and sodium ions are extracted from NaxCo1/3Ni1/3Mn1/3O2 oxides, resulting in a more complicated shape of the charge curves. It is of importance that, regardless of the nominal sodium content, the shape of charge curves for all oxides shows a tendency to become similar. The charge capacity corresponds to an extraction of more than 0.7 mol of alkaline ions, which exceeds significantly the amount of intercalated lithium ions. The high irreversible capacity is a specific feature of all sodium deficient transition metal oxides, used as cathodic materials in lithium ion batteries. However, the broad concentration range of stability of Co-containing oxides (i.e. 0.38 ≤ x ≤ 0.75) allows controlling the first irreversible capacity: Na-poor compositions with 0.38 ≤ x ≤0.5 exhibit a Coulombic efficiency of about 80-90 % depending on the charge/discharge rate, while the Columbic efficiency is varying between 40 and 60 % for Na-rich oxides with 0.67 ≤ x ≤0.75. This control of the first irreversible capacity cannot be achieved in case of NaxNi0.5Mn0.5O2 oxides due to their limited structural stability: 0.5≤x

Effect of Sodium Content on the Reversible Lithium Intercalation into

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