3O2 versus ... - ACS Publications

Research Article www.acsami.org

Layered P3‑NaxCo1/3Ni1/3Mn1/3O2 versus Spinel Li4Ti5O12 as a Positive and a Negative Electrode in a Full Sodium−Lithium Cell Svetlana Ivanova,† Ekaterina Zhecheva,† Rositsa Kukeva,† Diana Nihtianova,†,‡ Lyuben Mihaylov,§ Genoveva Atanasova,† and Radostina Stoyanova*,† †

Institute of General and Inorganic Chemistry and ‡Institute of Mineralogy and Crystallography, Bulgarian Academy of Sciences, 1113 Sofia, Bulgaria § Faculty of Chemistry and Pharmacy, Sofia University, 1164 Sofia, Bulgaria S Supporting Information *

ABSTRACT: The development of lithium and sodium ion batteries without using lithium and sodium metal as anodes gives the impetus for elaboration of low-cost and environmentally friendly energy storage devices. In this contribution we demonstrate the design and construction of a new type of hybrid sodium−lithium ion cell by using unique electrode combination (Li4Ti5O12 spinel as a negative electrode and layered Na3/4Co1/3Ni1/3Mn1/3O2 as a positive electrode) and conventional lithium electrolyte (LiPF6 salt dissolved in EC/DMC). The cell operates at an average potential of 2.35 V by delivering a reversible capacity of about 100 mAh/g. The mechanism of the electrochemical reaction in the full sodium− lithium ion cell is studied by means of postmortem analysis, as well as ex situ X-ray diffraction analysis, HR-TEM, and electron paramagnetic resonance spectroscopy (EPR). The changes in the surface composition of electrodes are examined by ex situ X-ray photoelectron spectroscopy (XPS). KEYWORDS: layered sodium transition metal oxides, lithium titanium spinel oxide, lithium ion batteries, sodium ion batteries, hybrid sodium−lithium ion batteries, TEM analysis, EPR spectroscopy, lithium and sodium intercalation

■

INTRODUCTION The manufacturing of the full electrochemical cell where lithium metal is not used as an anode gives the impetus for development of safe, low-cost, and environmentally friendly lithium ion batteries.1 The first commercial lithium ion battery comprised layered LiCoO2 as a positive electrode and graphite as a negative electrode.2 This battery is still used in mobile electronic devices, but its higher price and poor thermal stability are not in conformity with modern economic and social requirements. Nowadays, two groups of compounds deserve to be outlined as alternative positive and negative electrode materials. These are mixed lithium−transition metal oxides with a layered crystalline structure and a nominal composition of LiyCo1−2xNixMnxO2 (y ≥ 1, x = 1/3 and 1/2) and lithium titanium oxides with a spinel-type structure and a composition of Li 4 Ti 5 O 1 2 , respectively. 3 , 4 Layered LiyCo1−2xNixMnxO2 oxides are able to deliver above 4.0 V higher capacity at lower cost in comparison with that of the conventional LiCoO2 cathode.3,5 Spinel Li4Ti5O12 appears to be a safer negative electrode compared to the conventional carbon-based electrodes because it exhibits an operating voltage (i.e., 1.55 V versus Li/Li+) that is greater than the reduction potential of common electrolyte solvents.4,6 The main advantage of Li4Ti5O12 is its negligible volume change during lithium ion intercalation, a property contributing to the excellent long-cycling performance of batteries.6 The common © 2016 American Chemical Society

property of both layered and spinel oxides is their ability to intercalate lithium ions reversibly. By combining the intercalation properties of both oxides, a new full lithium cell with layered oxide LiCo1/3Ni1/3Mn1/3O2 and spinel Li4Ti5O12 as positive and negative electrodes is constructed so that a significant improvement in the cycling stability and rate performance is achieved versus its LiCoO2−graphite cell analogue.7,8 The improved performance of the full lithium cell is sufficient to compensate, to a great extent, for its main drawback related to its reduced energy density. Using the same concept of battery function, sodium ion batteries have recently been proposed as a low-cost alternative to the lithium ones.9 The main group of cathode materials for sodium ion cells is sodium−transition metal oxides with a layered structure as in the case of lithium−transition metal oxides.10,11 The bigger ionic radius of Na+ compared to that of Li+ gives rise to the different layer stacking and sodium site symmetry, as a result of which a variety of structures has been established for sodium−transition metal oxides.10,11 The crystalline structure is built up of discrete transition metal oxide layers, where sodium ions are sandwiched between the layers in such a way as to occupy either octahedral or prismatic Received: April 28, 2016 Accepted: June 17, 2016 Published: June 17, 2016 17321

DOI: 10.1021/acsami.6b05075 ACS Appl. Mater. Interfaces 2016, 8, 17321−17333

Research Article

ACS Applied Materials & Interfaces

tions are prepared by standard procedures described elsewhere.14,4 For the sake of comparison, both layered and spinel oxides were tested in lithium half-cells versus lithium metal anode. The mechanism of the electrochemical reaction in the full sodium−lithium ion cell was studied by ex situ X-ray diffraction analysis, HR-TEM, and electron paramagnetic resonance spectroscopy (EPR). The surface composition of both the negative and positive electrodes was probed by ex situ X-ray photoelectron spectroscopy (XPS).

sites. On the basis of 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 into O3-, P3-, and P2-types according to the notation, proposed by Delmas et al.12 Among different modifications, the most attractive structures, from the electrochemical point of view, are the P2- and P3-types, where the sodium intercalation proceeds at voltage higher than 3.5 V versus Na+/Na.10,11 Sodium titanium oxides are regarded as electrodes of choice for sodium ion batteries, where layered oxide Na2Ti3O7 is able to intercalate Na+ reversibly at the lowest voltage known for the oxide insertion electrode.13 The cell configuration includes layered P2-Na0.80Li0.12Ni0.22Mn0.66O2 and layered Na2Ti3O7 as positive and negative electrodes, which represents a successful approach to create a high-performance sodium ion battery.14 Recently, it has been demonstrated that sodium-deficient transition metal oxides with a P3-type of structure can intercalate not only Na+ ions, but also some Li+ ions.15−17 Although the layered structure is retained after Na + intercalation, the Li+ insertion causes a phase transformation from a P3- into O3-type of structure.18,19 The intercalation properties of sodium-deficient transition metal oxides allow using them as cathodes in lithium cells versus lithium metal anode.15−19 The so-designed “hybrid” lithium ion cell represents a low-cost alternative to the conventional lithium ion cells using lithium compounds. The most suitable compositions are NaxCo1−2yNiyMnyO2 (0.38 ≤ x ≤ 0.75 and y = 1/3; 0.50 ≤ x < 0.75 and y = 1/2) exhibiting a P3-type of structure.17,15 The lithium ion mobility inside the layered structure is enabled due to the electrochemical activity of all transition metal ions: the redox couples Ni2+/Ni3+,4+ and Co3+/ Co4+ that operate above 3.5 V, as well as the Mn3+/Mn4+ ion pair, which is active below 3.0 V.17−19 The reversible discharge capacity reaches up to 120 mAh/g when the cell is being cycled between 2.0 and 4.5 V.17−19 A specific feature to intercalate both Li+ and Na+ ions is also observed for spinel Li4Ti5O12.20,21 As a result, spinel Li4Ti5O12 serves as electrode in sodium cells versus sodium metal anode.20 The half-sodium cells deliver discharge capacity of about 155 mAh/g at an average voltage of 0.91 V.20 The ability of both layered NaxCo1−2yNiyMnyO2 and spinel Li4Ti5O12 oxides to intercalate Li+ and Na+ ions simultaneously provides an opportunity to design new types of hybrid alkaline ion batteries that combine the advantages of both lithium and sodium ion batteries. The approach based on the energy storage due to dual-metal ions intercalation seems very attractive in order to improve the electrochemical performance of batteries as well as to widen their applications.22 For example, the combination of the well-known LiMn2O4 spinel with Na0.22MnO2 oxide into a hybrid ion battery using brine as an electrolyte can be applied for electrochemical extraction and purification of lithium.23 In general, finding the efficient combination of materials for the assembly of a full electrochemical cell is the shortest route to design a new type of hybrid alkaline ion batteries. In this contribution, we demonstrate the elaboration of a new type of hybrid sodium−lithium ion cell by using unique electrode combination (Li4Ti5O12 spinel as a negative electrode and layered P3-Na3/4Co1/3Ni1/3Mn1/3O2 as a positive electrode) and conventional lithium electrolyte (LiPF6 salt dissolved in EC/DMC). The layered P3-Na3/4Co1/3Ni1/3Mn1/3O2 oxide was chosen in view of its higher reversible capacity among other sodium deficient oxides.14 Both layered and spinel composi-

■

EXPERIMENTAL SECTION

Synthesis of Oxides. The Li4Ti5O12 oxides were obtained by mixing Li acetate and TiO2 (anatase) in ethanol in an agate mortar for several minutes. The solid residual was heated at 400 °C, followed by thermal annealing at 750 °C for 12 h and 800 °C for 12 h. As a result, Li4Ti5O12 with a lattice parameter of a = 8.3585 Å is obtained (See Figure 1a and Supporting Information).

Figure 1. Charging/discharging curves for Li4Ti5O12 spinel in a halfcell after 1, 2, 5, and 10 cycles. The electrochemical test is started with discharging mode. Inset: the capacity stability curves, where the open and full symbols correspond to charging and discharging capacity, respectively. The Na3/4Co1/3Ni1/3Mn1/3O2 oxide was 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 was fixed at Na/Co/ Ni/Mn = 3/4:1/3:1/3:1/3. The solid residual was heated at 400 °C, followed by thermal annealing at 800 °C for 10 h. This method allows us to obtain Na3/4Co1/3Ni1/3Mn1/3O2 with a P3-type of structure and lattice parameters of a = 2.8413 Å and c = 16.6777 Å (Figure 1b, Supporting Information). In addition, a mixture of NiO-like phase is also obtained. These results are in agreement with our previous structure data on NaxCo1/3Ni1/3Mn1/3O2.17 Structure and Morphology Characterization. The X-ray structural analysis was made on Bruker Advance D8 powder diffractometer with CuKα radiation. Step-scan recordings for structural refinement by Rietveld’s method were carried out using 2θ steps of 0.02° 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 half-width parameters were determined. The computer FullProf Suite Program (1.00) was used in the calculations.24 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 min. The suspensions were dripped on standard holey carbon/Cu grids. The analysis was carried out by the Digital Micrograph software. 17322


Research Article

ACS Applied Materials & Interfaces 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 120−450 K temperature interval. Surface Analysis. X-ray photoelectron spectroscopy (XPS) was carried out using AXIS Supra electron-spectrometer (Kratos Analitycal Ltd.) with monochromatic AlKα source (excitation energy of 1486.6 eV) and charge neutralization system. The binding energies (BE) were determined with an accuracy of ±0.1 eV. The processing of the recorded XPS spectra was performed, using a symmetrical Gaussian− Lorentzian curve fitting after Shirley’s type subtraction of the background. The commercial data-processing software of Kratos Analytical Ltd. was applied for the calculation of element concentrations in atomic percent (atom. %). Electrochemical Characterization. The electrochemical charge− discharge of Na3/4Co1/3Ni1/3Mn1/3O2 and Li4Ti5O12 was carried out by using two-electrode cells versus lithium anode and LiPF6 (EC/DMC) electrolyte. The positive electrode, supported on an aluminum foil, was a mixture containing 80% of the active composition Na3/4Co1/3Ni1/3Mn1/3O2, 7.5% KS 6L graphite (TIMCAL), 7.5% Super C65 (TIMCAL), and 5% polyvinylidene fluoride (PVDF). The electrolyte was a 1 M 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 a diameter of 18 mm. The full sodium−lithium cell is constructed by using two-electrode cells of the type Na x Co 1/3 Ni 1/3 Mn 1/3 O 2 |LiPF 6 (EC/DMC)| Li4/3Ti5/3O4. The specific capacity is calculated on the basis of the positive electrode. The electrochemical reactions were carried out using an eightchannel 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 ion per formula unit at the applied current intensity. The cells were mounted in a drybox under Ar atmosphere. The half cells were cycled between 4.4 and 1.8 V for Na3/4Co1/3Ni1/3Mn1/3O2 and between 2.5 and 1.0 V for Li4Ti5O12 at C/20 rate. The potential window for the full cell is 1.0−3.5 V at a rate of C/20. The structural changes in the electrode compositions during reversible intercalation and deintercalation were analyzed with full sodium−lithium cells stopped at selected potentials. The electrochemical cells were disassembled insight a glovebox, followed by removal and washing of the working electrodes with EC. The electrode samples were covered with parafilm for the XRD experiments in order to avoid any water contamination. The specimens were dispersed in acetone for the TEM experiments, and then the suspensions were dripped on standard holey carbon/Cu grids.

summary, the electrochemical half-cell reaction can be written as Li4/3Ti5/3O4 ↔ Li 7/3Ti5/3O4 + Li+ + e−

Because of the high oxidation state of Ti ions in Li4/3Ti5/3O4 (i.e., Ti4+), the electrochemical cycling test starts by a discharge mode only, i.e., by a lithium intercalation into Li4/3Ti5/3O4. In comparison with Li 4 / 3 Ti 5 / 3 O 4 spinel, layered P3Na3/4Co1/3Ni1/3Mn1/3O2 contains transition metal ions in mixed oxidation states that allow the cell to start either by a discharging or a charging mode. Our previous studies show that during the first discharge Li+ ions are intercalated into NaxCo1/3Ni1/3Mn1/3O2 in an amount compensating the sodium deficiency.17 Lithium intercalation takes place due to the electrochemical activity of all transition metal ions. To complete the full cell, there is a need to analyze the intercalation properties of P3-Na3/4Co1/3Ni1/3Mn1/3O2 in a lithium half-cell starting with a charging mode, where Na ions extraction from the layered structure is expected. Figure 2 compares the charge/discharge curves of P3Na3/4Co1/3Ni1/3Mn1/3O2 in half cells versus lithium anode,

Figure 2. Charging/discharging curves for P3Na0.75Co1/3Ni1/3Mn1/3O2 in a half-cell after 1, 2, and 10 cycles. The electrochemical test is started with a discharging (a) and charging mode (b). (c) First (1) and second (2) charge/discharge curves expressed as derivatives (dQ/dV) for P3-Na0.75Co1/3Ni1/3Mn1/3O2 versus Li anode in a lithium half-cell. (d) Capacity stability curves for P3-Na0.75Co1/3Ni1/3Mn1/3O2 when the test starts with charging (ch) and discharging mode (dish). Full and open symbols correspond to a discharging and charging capacity.

■

RESULTS AND DISCUSSIONS Electrochemical Properties of Layered P3Na3/4Co1/3Ni1/3Mn1/3O2 and Li4Ti5O12 Spinel in Half Cells versus Li Anode. Prior to the electrochemical test in a full cell, both Li4Ti5O12 and Na3/4Co1/3Ni1/3Mn1/3O2 oxides were examined in half cells versus lithium anode. Figure 1 shows the charge/discharge curves for Li4Ti5O12 spinel in a lithium halfcell. To rationalize the electrochemical intercalation properties of Li4Ti5O12, the composition is expressed by a spinel formula unit Li4/3Ti5/3O4. The spinel oxide, obtained by us, displays the typical electrochemical behavior:4 Charge/discharge processes are taking place at a voltage plateau of 1.57/1.56 V due to the Ti3+/Ti4+ redox couple. The reversible capacity reaches 165 mAh/g, which corresponds to 1 mol of inserted Li+ per formula unit Li4/3Ti5/3O4. After first few cycles, the cycling stability of Li4/3Ti5/3O4 becomes very good with a coulombic efficiency close to 100%. The occurrence of a voltage plateau at 1.55 V reflects a phase transformation from the spinel Li4/3Ti5/3O4 into a rocksalt Li7/3Ti5/3O4 during lithium intercalation.4 In

which are started by a discharge and by a charge mode. When the cell is first discharged, the oxide delivers a specific capacity of 66 mAh/g corresponding to an insertion of about 1/5 mol of Li+ into P3-Na3/4Co1/3Ni1/3Mn1/3O2 (Figure 2a). In accordance with our previous data,17 the lithium insertion is accomplished passing through two voltage plateaus due to a selective reduction of transition metal ions: Between 3.7 and 3.5 V there is a reduction of highly oxidized Ni3+/4+ into Ni2+, whereas below 3.0 V, reduction of Mn4+ into Mn3+ is occurring together with a partial reduction of Co3+ to Co2+. The reverse process of charge leads to a release of a higher capacity (165 mAh/g), which indicates that both lithium and sodium ions are extracted (1/2 mol of alkaline ions). In addition to Ni2+/Ni3+,4+ and Mn3+/Mn4+ couples, the Co3+/Co4+ pair also becomes 17323


Research Article

ACS Applied Materials & Interfaces active above 3.5 V. Upon prolonged cycling, the charge/ discharge curves become less structured. This process has been shown to originate from continuous changes in structure and chemical composition of P3-NaxCo1/3Ni1/3Mn1/3O2 oxides (0.38 ≤ x ≤ 0.75), culminating after several cycles in the formation of mixed layered oxides having an O3-type of structure, where both lithium and sodium ions are appearing simultaneously in the alkaline layers: Liy→1Nax∼0.03Co1/3Ni1/3Mn1/3O2.17 The reversible electrochemical reaction proceeds, furthermore, owing to the lithium intercalation in the framework of the O3-type of structure.17 When the cell starts with a charging mode, a specific capacity of 80 mAh/g is achieved, which is related to extraction of 1/4 mol of Na+ from P3-Na3/4Co1/3Ni1/3Mn1/3O2 (Figure 2b). The continuous changing in the charge curve shape indicates that the sodium extraction is taking place in the framework of the single layered structure. This process can be, most probably, associated with redox properties of the nickel ions. The reverse process of discharging occurs by delivering a higher capacity (i.e., 125 mAh/g, Figure 2b). This reveals that 2/5 mol of Li+ and Na+ ions are inserted into the oxide. As in the case when the experiment is starting with a discharging mode, two wellpronounced regions above and below 3.0 V are clearly distinguished (Figure 2c, where the curves are expressed as first derivatives). The two voltage regions can mainly be associated with Ni2+/Ni3+,4+ and Mn3+/Mn4+ couples as was previously observed. The characteristic feature of sodium deficient oxides, used as cathodes in lithium half cells, is their irreversible capacity. It is worth mentioning that when the cell starts with a charging mode the irreversibility drastically decreases from 90 to 50 mAh/g (Figure 2a,b). Irrespective of the starting mode of charging or discharging, the shape of electrochemical curves after the first cycle becomes quite similar, which is indicative that the same electrochemical mechanism is operating (Figure 2). In comparison with the cell starting with a discharging mode, the cell beginning with a charging mode delivers slightly higher reversible capacity: 110 versus 100 mAh/g, respectively. Both cells show a coulombic efficiency higher than 99% after the first 6 cycles. Starting with a charging mode, the electrochemical half-cell reaction can be written as

Figure 3. First (1) and second (2) charge/discharge curves expressed as derivatives (dQ/dV) for P3-Na0.75Co1/3Ni1/3Mn1/3O2 versus Li4/3Ti5/3O4 in a full sodium−lithium cell (a). Charge/discharge curves (b) and corresponding capacity stability curves (c). The open and full symbols correspond to discharging and charging capacity, respectively.

electrochemical properties, all curves are represented as first derivatives. In addition, the charging/discharging curves as first derivatives for P3-Na3/4Co1/3Ni1/3Mn1/3O2 and Li4/3Ti5/3O4 in lithium half-cells are also given. During the first charging of the full cell, three well-defined peaks can be distinguished at 2.39, 2.60, and 2.87 V. Although the three voltage peaks fall into the range of the potential difference between individual P3Na3/4Co1/3Ni1/3Mn1/3O2 and Li4/3Ti5/3O4 oxides, the first charging profile of the full cell is more structured in comparison with those for lithium half-cells. This indicates that the function of the full cell is not a simple sum of the functions of the halfcells. It is of importance that the delivered capacity reaches a value of 135 ± 5 mAh/g, which matches the mass balance of the two electrodes and their individual capacities in lithium half-cells (Figures 1 and 2). During the reverse process of discharging, the electrochemical curve displays three well-resolved peaks at 2.47, 2.25, and 1.83 V, respectively. The three voltage peaks can also be distinguished in t he discha rging curve for P 3Na3/4Co1/3Ni1/3Mn1/3O2 versus Li anode, i.e., 4.02, 3.75, and 3.41 V vs Li (Figure 2c). It is noticeable that all the voltage peaks above 3.0 V are associated with electrochemical activity of nickel and cobalt ions only. In addition, close inspection of the discharging profiles reveals that the voltage peak below 3.0 V (i.e., at 2.9 V) for the half-cell is not active in the full cell (Figures 3a and 2c). This voltage peak is related mainly to a Mn3+/Mn4+ redox couple. The comparison of the discharging profiles for the half-cells and the full cells reveals that the electrochemical response of the full cell is mainly due to the

P 3‐Na3/4Co1/3Ni1/3Mn1/3O2 → O3‐Li yNax ∼ 0.03Co1/3Ni1/3Mn1/3O2 ↔ O3‐Li y − nNax ∼ 0.03Co1/3Ni1/3Mn1/3O2 + n Li+ + ne−

Full Sodium−Lithium Cell Combining Layered P3Na3/4Co1/3Ni1/3Mn1/3O2 and Li4Ti5O12 Spinel Electrodes. The electrochemical behavior of layered P3Na3/4Co1/3Ni1/3Mn1/3O2 and Li4/3Ti5/3O4 spinel in lithium half-cells allow us to combine them as positive and negative electrodes in a full sodium−lithium cell. The mass balance between two electrode materials plays an important role in designing the full cell. Our results show that the best electrochemical response is achieved when the molar ratio of the positive electrode to the negative one is 2.0−1.0 and the open circuit voltage is 0.25 V. Therefore, the cycling test starts with a charging mode in respect of the positive electrode, where Na is expected to be extracted from P3Na3/4Co1/3Ni1/3Mn1/3O2. Figure 3 shows the first and second charging−discharging profiles. For better comparison of the 17324


Research Article

ACS Applied Materials & Interfaces redox Ni3+,4+/Ni2+ and Co3+/Co4+ ion pairs. The discharge capacity delivered by the full cell between 3.0 and 1.0 V amounts to 100 ± 5 mAh/g. During the second cycle, both charging and discharging curves become less structured, and they consist of broad peak with center of gravity of 2.38/2.32 V, respectively (Figure 3b). The same picture is observed with P3-Na3/4Co1/3Ni1/3Mn1/3O2 in a lithium half-cell: The broad oxidation and reduction peaks at 3.98 and 3.97 V vs Li appear, which are related to redox couples Ni3+,4+/Ni2+ and Co3+/Co4+. In the voltage region below 3.2 V vs Li, P3-Na3/4Co1/3Ni1/3Mn1/3O2 additionally displays two reduction and oxidation peaks at 2.79 and 3.10 V due to the redox ion pair Mn3+/Mn4+. These peaks are again not observed in the charging/discharging curves for P3Na3/4Co1/3Ni1/3Mn1/3O2 in the full cell vs Li4/3Ti5/3O4. This indicates that the voltage peaks in the full cell originate from the redox couples Ni3+,4+/Ni2+ and Co3+/Co4+ only. The full cell comprising Li4/3Ti5/3O4/LiPF6(EC/DMC)/P3Na3/4Co1/3Ni1/3Mn1/3O2 delivers during prolonged cycling a reversible capacity at an average potential of 2.32−2.38 V (Figure 3c,d). The capacity fade can be related to a restricted supply of movable alkaline ions provided by the layered oxide, as well as an increase in resistance of components by formation of solid electrolyte interface (SEI, see the XPS data). Mechanism of the Electrochemical Reaction of the Full Sodium−Lithium Cell. Postmortem Analysis. To understand the composition changes occurring in the negative and the positive electrodes during full cell operation, a postmortem analysis was undertaken. The full cell is first charged up to 3.5 V. Then, the cell is disassembled, and the positive and negative electrodes are tested in the half-cells versus Li anode. The corresponding charging/discharging curves are given in Figure 4. In the case of the positive

charged up to 3.5 V that 2/3 mol of alkaline ions are inserted into Li4/3Ti5/3O4. The next charging curve for postmortem Li4/3Ti5/3O4 versus Li anode is also identical with that of the fresh Li4/3Ti5/3O4. These data supply clear evidence for the participation of the Li4/3Ti5/3O4 spinel in the electrochemical process of the full cell operation. However, it is not clear whether Li+ and/or Na+ ions are being inserted into Li4/3Ti5/3O4. The postmortem analysis was also carried out for the positive and negative electrodes when the full sodium−lithium cell is undergoing the first complete cycle of charging up to 3.5 V and discharging down to 1.0 V (Figure 4). In a lithium half-cell, the charge of postmortem-derived oxide P3Na3/4Co1/3Ni1/3Mn1/3O2 displays an extraction of 0.15 mol of alkaline ions, whereupon the potential of extraction remains unchanged. Complementary to the positive electrode, the postmortem-derived oxide Li4/3Ti5/3O4 intercalates Li+ ions in the same amount as that in the case of the fresh spinel at the same potential of 1.55 V. Taking into account the data from postmortem analyses, the electrochemical process of charging of the full sodium−lithium cell can be described as an extraction of 1/3 mol Na+ ions from the positive electrode P3-Na3/4Co1/3Ni1/3Mn1/3O2 in parallel to insertion of 2/3 mol alkaline ions into negative electrode Li4/3Ti5/3O4. During the reverse process of discharging, the opposite reactions of insertion and extraction of alkaline ions are occurring. Li4Ti5O12 Used as a Negative Electrode in a Full Sodium−Lithium Cell. Structure Changes in Li4Ti5O12 during the Charging/Discharging Process. Further light is thrown on the electrochemical process by ex situ XRD, TEM, EPR, and XPS measurements of both positive and negative electrodes. Three types of electrode samples were studied: P3Na3/4Co1/3Ni1/3Mn1/3O2 and Li4/3Ti5/3O4 after the first charging up to 3.5 V, P3-Na3/4Co1/3Ni1/3Mn1/3O2 and Li4/3Ti5/3O4 after the first complete cycle down to 1.0 V, and P3-Na3/4Co1/3Ni1/3Mn1/3O2 and Li4/3Ti5/3O4 after the 10 cycles between 1.0 and 3.5 V and the cell stopped at 1.0 V. Figure 5 shows the XRD patterns of Li4/3Ti5/3O4 electrode compositions. After the first charging of the full cell up to 3.5 V, the XRD pattern consists of a mixture of two cubic phases. The first cubic phase has a lattice parameter slightly higher than that of the pristine composition, i.e., a = 8.3716 Å vs a = 8.3585 Å, whereas a greater lattice parameter is observed for the second phase, i.e., a = 8.4087 Å. The appearance of two phases is

Figure 4. First charge and discharge curves in a lithium half-cell for NaxCo1/3Ni1/3Mn1/3O2 (A) and Li4/3Ti5/3O4 (B) obtained after the first charge of the full cell up to 3.5 V (a) and after the further discharge down to 1.0 V (b).

electrode, the electrochemical test starts with a charging mode, i.e., during which alkaline ions are extracted (Figure 4A), whereas for the negative electrode, the cell is discharging, i.e., during its course alkaline ions are inserted (Figure 4B). Although for the fresh electrode 0.32 mol of Na+ is extracted after the first charging in a lithium half-cell, for the postmortem positive electrode, it is possible to extract no more than 0.02 mol of Na+. The next discharging curves for both fresh and postmortem electrodes coincide in terms of insertion potential and capacity. This reveals that during the first charging of the full sodium−lithium cell almost all of the moveable Na+ ions are extracted from P3-Na3/4Co1/3Ni1/3Mn1/3O2 whereupon the layered structure is preserved. In the case of the negative electrode, it is possible to insert only 1/3 mol of Li+ into postmortem Li4/3Ti5/3O4 at a potential matching that of the fresh Li4/3Ti5/3O4 (i.e., at 1.56 V). This means that in a full cell

Figure 5. XRD patterns of pristine (a) Li4/3Ti5/3O4 and (b) Li4/3Ti5/3O4 electrodes delivering a capacity of 145 mAh/g after the first charging to 3.5 V. (c) Li4/3Ti5/3O4 delivering a capacity of 145 and 99 mAh/g after the first charging to 3.5 V followed by discharging to 1.0 V and (d) Li4/3Ti5/3O4 after 10 cycles between 1.0 and 3.5 V with reversible capacity of 66 mAh/g (stopped at 1.0 V). All Li4/3Ti5/3O4 electrodes are tested in a full sodium−lithium cell. 17325


Research Article


tively.20 Computational results point out that the lattice parameters of the solid solutions between Li7/3Ti5/3O4 and Na6/3Li1/3Ti5/3O4 obey Vegard’s law.20 On the basis of the correlation between the structure and composition between Li 4/3 Ti 5/3 O 4 , Li 7/3 Ti 5/3 O 4 and Na6/3Li1/3Ti5/3O4, we can rationalize the ex situ XRD data on Li4/3Ti5/3O4 negative electrode used in a full sodium−lithium cell. It appears that the first cubic phase has a slightly higher lattice parameter than that of the pristine Li4/3Ti5/3O4, which is due to a lithium-intercalated phase, i.e., Li7/3Ti5/3O4, whereas the second cubic phase with higher lattice parameter originates from sodium-intercalated phase, i.e., Li(7−x)/3Nax/3Ti5/3O4. Taking into account the previously calculated concentration dependence of the lattice parameter,20 it seems that at least 1/3 mol Na+ is present in the cubic phase. The XRD data give clear evidence that during the charging of the full cell up to 3.5 V both Na+ and Li+ ions are intercalated into Li4/3Ti5/3O4. This is also consistent with the shape of the charging curve, which seems to be more “structured” in comparison with those for the lithium half-cells (Figures 1−3). The amount of Na in the cubic phase is sensitive toward the process of charging and discharging. During the first charging of the full cell, both Li+ and Na+ are being intercalated into Li4/3Ti5/3O4 spinel, with the ratio of Na-containing phase being higher. After the reverse process of charging, there is a simultaneous extraction of Li+ and Na+ ions leading to a restoration of Li4/3Ti5/3O4 spinel phase. However, some sodium-containing cubic phase coexists with Li4/3Ti5/3O4. During prolonged cycling, the sodium phase becomes enriched in sodium, but its amount decreases. This means that the Li+ ions are the main intercalation species during cell operation. The coexistence of two lithium- and sodium-intercalated phases is also proved by ex situ TEM analysis. Figure 6 gives the bright-field, SAED, and HRTEM images of cycled composition Li4/3Ti5/3O4 after 10 cycles between 1.0 and 3.5 V (the cell test is stopped at 1.0 V, i.e., the oxide is in a discharged state). The cycled oxide consists of well-faceted particles with dimensions varying around 500−700 nm. The lattice parameter estimated by SAED, taken along the [0 1 1] direction, is a = 8.43 Å, which coincides with that determined by ex situ XRD analysis of Na intercalated Li4/3Ti5/3O4 (a = 8.4359 Å). The HRTEM images of selected particles show an occurrence of closely interconnected nanodomains with different lattice parameters. On the basis of the calculations along [011] direction, it is possible to distinguish domains with at least four different lattice parameters: in addition to domains with a = 0.838 and 0.843 nm corresponding to lithium- and sodium-intercalated Li4/3Ti5/3O4, i.e., as it was observed by XRD, there are fingerprints for domains with a = 0.836 and 0.872 nm related to pristine Li 4 / 3 Ti 5 / 3 O 4 and to sodium pure phase Na6/3Li1/3Ti5/3O4. This result is a clear evidence for a formation of both lithium- and sodium-intercalated phases during prolonged cycling of Li4/3Ti5/3O4 in a sodium−lithium full cell. The coexistence of domains with slightly different unit cell parameters is consistent with the observed broadening of XRD diffraction lines for cycled electrodes (Figure 5). It is interesting that the domain formation inside a single particle has already been observed by STEM for Li4/3Ti5/3O4 cycled in a half-cell versus Na anode.20 The proposed mechanism includes simultaneous generation of both Na- and Li-inserted phases, i.e., Na6LiTi5O12 and Li7Ti5O12, which are completely transformed into the pristine phase Li4Ti5O12 after the reverse process of Na+ extraction.20 Contrary to this

clearly demonstrated by a deconvolution of the (111) diffraction peak (Figure 5). It is noticeable that both diffraction peaks become broadened in comparison with that of pristine Li4/3Ti5/3O4. The intensity of the phase with a larger lattice parameter is twice as high. The reverse process of discharging causes further changes in phase composition and structure (Figure 5). The pristine spinel phase Li4/3Ti5/3O4 is restored with the exception of the line widths of diffraction peaks. The lattice parameter of the second cubic phase decreases slightly and it remains higher in comparison with that of the pristine composition (a = 8.3886 Å, Figure 5). The two cubic phases coexist in nearly equal amounts. After prolonged cycling, there is a clear differentiation between the two cubic phases with respect to the lattice parameter (Figure 5). The lattice parameter of the first cubic phase is again slightly higher than that of the pristine Li4/3Ti5/3O4 (a = 8.3804 Å), whereas the second cubic phase displays a highest lattice parameter and its amount decreases (a = 8.4359 Å). The line widths of all diffraction peaks do not undergo any measurable changes. The differentiation between the cubic phases on the basis of the lattice parameter manifests the changes in the composition of Li4/3Ti5/3O4 occurring during the electrochemical reaction. Li4/3Ti5/3O4 exhibits a spinel type of structure, where 1 Li+ and 5 Ti4+ ions occupy the octahedral 16d sites, whereas the rest 3 Li+ ions occupy tetrahedral 8a positions.4 In a lithium half-cell, Li4/3Ti5/3O4 is able to intercalate 1 Li+ ions into vacant 16c spinel sites, which in their turn provoke a collective transition of all Li+ ions from tetrahedral 8a sites into the octahedral 16c sites.4,6 This is a reversible two-phase reaction between rocksalt and spinel types of structure. The characteristic feature of Li4/3Ti5/3O4 spinel is a negligibly small difference in the lattice parameters of the pristine and intercalated phase. The lattice parameter of Li4/3Ti5/3O4 has been reported to be a = 8.36 Å, whereas in the case of Li7/3Ti5/3O4, the lattice parameter is 8.37 Å.6 Recently, a new intercalation property for Li4/3Ti5/3O4 spinel has been reported. In a sodium half-cell, it has been demonstrated that Li4/3Ti5/3O4 spinel has intercalated sodium ions at average voltage of 0.91 V vs Na+/Na.20,21 The electrochemical reaction proceeds via a three-phase separation mechanism according to the equation20 2Li4/3Ti5/3O4 + 6/3Na + + 6e− ↔ Li 7/3Ti5/3O4 + Na6/3Li1/3Ti5/3O4

As in the case of Li+, Na+ ions occupy the vacant 16c spinel sites leading to the formation of Na6/3Li1/3Ti5/3O4 with a rocksalt type of structure. Contrary to lithium-intercalated Li4/3Ti5/3O4, the sodium-intercalated Li4/3Ti5/3O4 displays a much greater lattice volume expansion. The lattice parameter of Na6/3Li1/3Ti5/3O4 is of 8.72 Å, which means a lattice volume expansion of 12.5%.20 According to DFT-based first-principle calculations, it has been found out that the formation of solid solutions between Li4/3Ti5/3O4, Li7/3Ti5/3O4 and Na6/3Li1/3Ti5/3O4 is an energetically unfavorable process.20 However, the in situ synchrotron XRD measurements of the insertion of Na+ ions into Li4/3Ti5/3O4 show the formation of solid solutions only in a limited concentration range, which are near to the stoichiometric end members, i.e., close to Li(4+x)/3Nax/3Ti5/3O4, Li(7−x)/3Nax/3Ti5/3O4, and to Na(6−x)/3Lix/3Ti5/3O3, respec17326


Research Article


Figure 7. BF-STEM images and corresponding composition map of TiKα 1 , O Kα 1 , and Na Kα 1 − 2 for cycled electrodes NaxCo1/3Ni1/3Mn1/3O2 with nominal sodium content of x = 0.75.

Figure 6. Bright-field images, SAED (top), and HRTEM (bottom) of cycled composition Li4/3Ti5/3O4 after 10 cycles between 1.0 and 3.5 V (the cell test is stopped at 1.0 V, i.e., the oxide is in a charged state).

as well as elements P and F belonging to LiPF6 salt. The information on the chemical state of these elements is of importance to rationalize not only the interaction between Li4Ti5O12 with alkaline ions but also the interaction between electrode and electrolyte. In a lithium half-cell, Li4Ti5O12 operates at a potential, which is above the potential for the formation of SEI due to the reduction of electrolyte solution.4 However, a SEI image of the surface of Li4Ti5O12 has experimentally been observed, which is supposed to be formed due to the presence of conducting carbon in the electrode acting as active sites.26 For this reason, XPS is applied to monitor the surface of Li4Ti5O12 electrode working in the lithium−sodium full cell. In the energy region of Li 1s, the XPS spectrum of pristine Li4Ti5O12 electrode displays two overlapping signals with binding energies of 56.1 and 52.5 eV (Figure 8). The lowenergy signal originates from lithium in the spinel oxide as was previously observed. The high-energy signal at 56.1 eV can be associated with fluoride ions in LiF characterized by a binding energy of 55.6 ± 0.2 eV.27 The occurrence of LiF can be explained in terms of the surface interaction between Li4Ti5O12 spinel and PVDF binder used for the manufacturing of the pristine electrode. The surface interaction is also observed in the energy region of F 1s, where two overlapping signals with binding energies of 687.5 and 685.2 eV are clearly distinguished. The high-energy signal originates from the fluorine atom in PVDF binder,28 whereas the low-energy signal, i.e., 685.2 eV, corresponds to fluoride ion in LiF. This outlines the fact that prior to the electrochemical reaction the Li4Ti5O12 spinels exhibit a surface reactivity toward PVDF binder. It is interesting that Li4Ti5O12 spinel has been shown to display better electrochemical properties (in terms of specific capacity and cycling stability) when carboxy methyl cellulose (CMC) is used as binder instead of PVDF.29 However, following the main purpose of this study, we used only PVDF as a binder.

mechanism, a Na-substituted phase Li4−xNaxTi5O12 (0 < x < 3) coexisting with Li4Ti5O12 has been identified during numerous sodium insertion−extraction cycles.23 It has been supposed that both Na- and Li-containing phases are involved in the charge− discharge reactions.25 The Na-inserted phase Li4−xNaxTi5O12 (0 < x < 3) has an almost constant lattice parameter of about 0.845−0.846 nm.25 The comparison shows an amazing consistency between the lattice parameters of the “extra” Nainserted phase identified by Kitta et al.25 and Na-containing phase determined by us after 10 cycles between 1.0 and 3.5 V of the full cell. This confirms once again that in the full cell where lithium electrolyte is used Na+ intercalation into negative electrode Li4Ti5O12 takes place as in the case of a sodium halfcell where sodium electrolyte is ised.20,25 Using BF-STEM, the composition of the cycled phase is evaluated after prolonged cycling (Figure 7). The results show a homogeneous distribution of the Ti and O elements in particles, therefore giving evidence for a stability of the cubic matrix during alkaline ions intercalation. In addition, the Na element is also detected, and its distribution follows the profile of the Ti one. The sodium amount is nearly 10 times lower than that of the Ti element, which is further proof for Na intercalation into the spinel structure. However, the lower sodium content shows that Li+ ions are the main intercalation species. Surface Changes in Li4Ti5O12 during the Charging/ Discharging Process. The surface composition of the electrode is another parameter that affects the cell performance. In this study, the surface composition of Li4Ti5O12 electrodes operating in a full sodium−lithium cell is probed by XPS (Figure 8). For the sake of comparison, the pristine electrode Li4Ti5O12 is also shown in Figure 8. The analysis is focused on the energy regions where alkaline elements Li and Na appear, 17327


Research Article


with binding energies of 1071.9 and 1069.9 eV, respectively. The binding energy of the high-energy signal coincides with that characteristic of Na+ in NaF,30 therefore indicating a surface interaction of Li4Ti5O12 with Na+ ions. The occurrence of Na+ ions in the lithium electrolyte can be described as a consequence of the extraction of Na+ from Na3/4Co1/3Ni1/3Mn1/3O2 during the first charging of the full cell. This means that the surface deposition of NaF on the cubic oxide is taking place at the beginning of the electrochemical reaction due to electrolyte−electrode interaction. The signal with a lower binding energy falls within the range typical of Na+ ions bonded to oxygen in transition metal oxides. For example, the binding energy of Na+ in sodium cobalt−nickel−manganese oxides Na0.67Co1/3Mn1/3Ni1/3O2 with a P3-type of structure is 1070.6 eV.17 This allows assignment of the Na 1s signal at 1069.9 eV to Na+ ions bonded to oxygen inside the spinel structure. The XPS observation of the surface Na-containing cubic phase is also in agreement with XRD data (Figure 5), where bulk Na-containing phase is detected. Because Na 1s signal comes from only few top surface layers, the coexistence of Na-containing oxide and NaF implies that there is nonhomogeneous covering of the surface of the spinel particle. The amount of Na+ in Na-containing oxide is higher than that in surface located NaF, which is indicative of a favorable interaction of Na+ with the spinel phase instead of the electrolyte−electrode interaction. To assess the surface depth where sodium species are to be found, we also analyzed the Na 2s XPS spectrum (Figure 8). It has to be recalled that the Na 2s signal can be regarded as more “bulk-phase-like” in comparison with that of Na 1s. Moreover, the surface sensitivity toward Na 2s and Li 1s is similar because the electron energies for Na 2s and Li 1s fall within the same range. As in the case of Li 1s, two overlapping signals at 63.6 and 61.2 eV gives rise to the shape of the Na 2s spectrum. The signal with a binding energy of 63.6 eV is related to Na+ in NaF, i.e., about 63 eV for Na 2s.30 It appears that LiF and NaF are close to each other and that they are penetrating deeply inside of the oxide particle. The origin of the low-energy Na 2s signal can be attributed to Na+ ions bonded to oxygen in the cubic phase, as observed in the Na 1s region. However, the interpretation of the low-energy Na 2s signal is not straightforward because it is clearly affected by the signal due to Ti 3s. The P 2p spectrum of the Li4Ti5O12 electrode, after the first charging of the full sodium−lithium cell up to 3.5 V, consists of two broad signals with binding energies of 137.5 and 135.3 eV, respectively (Figure 8). According to the literature data, the peak having the lower binding energy, i.e., 134.5 eV is associated with LixPFy.31−34 Taking into account all of these data, one can identify LixPFy and/or NaxPFy compounds on the surface of the Li4Ti5O12 electrode. It is noticeable that mixed alkaline phospho-fluorine compounds are being formed during the first charging of the full cell up to 3.5 V. The reverse process of discharging of the full cell leads to increase in the intensities of the all signals due to the surface LixPFy and/or NaxPFy compounds, especially in the case of the Na 1s signal at 1071.2 eV. This means that during the cell operation there is a continuous growth of the surface layer. The observation of the Li 1s and Na 1s signals due to Li- and Nacontaining cubic phase in addition to the LixPFy and/or NaxPFy

Figure 8. XPS spectra in the Na 1s, Na 2s, Li 1s, F 1s, and P 2p regions for (a) pristine Li4Ti5O12 and (b) Li4Ti5O12 electrodes after the first charging of the full cell to 3.5 V, (c) after the first cycle of the full cell to 1.0 V, and (d) after 10 cycles between 1.0−3.5 V (the cell is stopped at 1.0 V). The Ti LMM, Ti 3s, and Cu LMM (due to copper support) signals are also indicated in the Na 1s, Na 2s, and F 1s regions.

In general, lithium on the surface of the pristine Li4Ti5O12 electrode is in the form of PVDF-generated LiF and unreacted Li-containing spinel phase. It is noticeable that the pristine electrode composition does not show any XPS signals due to Na and P elements (Figure 8). However, it has to be taken into account that the sodium signals of Na 1s and Na 2s are affected by the presence of the main titanium element due to the their overlapping with Ti LMM and Ti 3s, respectively. The comparison shows that the Na 1s signal is less affected (Figure 8). After the first charging of the full cell up to 3.5 V, the Li 1s and F 1s spectra retain their complex profiles without undergoing any alteration in the binding energies. This means that the surface forms of Li are preserved. The only parameter that is changed is the ratio between them. The signal intensities of Li 1s and F 1s come from LiF increase (Figure 8). This fact can be explained by “electrode−electrolyte” interaction stimulated during the electrochemical process. Contrary to Li, new signals grow in intensity in the energy region for Na 1s, Na 2s, and P 2p, respectively (Figure 8). The Na 1s spectrum can be deconvoluted into two components 17328


Research Article


4.81 Å being higher (Figure 9). This means that the prolonged cycling of the full cell causes a continuous conversion of the Na-containing phases into Li-containing ones. However, Na+ intercalation into layered oxide is also taking place. It is worth mentioning that in a lithium half-cell the electrochemical reaction comprises the intercalation of Li+ ions only.17 The ex situ TEM analysis provides further evidence for Na+ and Li+ intercalation into the positive electrode oxide. Figure 10

compounds is indication of nonhomogeneous growth of the surface film. The prolonged cycling of Li4Ti5O12 electrode causes further changes in the relative intensities of all the XPS signals; the XPS profiles for Na 1s, Li 1s, F 1s, and P 2p spectra are unaltered (Figure 8). This outlines the impact of electrochemical cycling on the thickness of the surface layer without changing the film composition. Layered P3-Na3/4Co1/3Ni1/3Mn1/3O2 Used as a Positive Electrode in a Full Sodium−Lithium Cell. Structure Changes in Layered P3-Na3/4Co1/3Ni1/3Mn1/3O2 during Charging/Discharging. In synchrony with Li4Ti5O12 negative electrode, layered P3-Na3/4Co1/3Ni1/3Mn1/3O2 positive electrode is also undergoing changes in its composition and structure (Figure 9). Because of a strong broadening in all

Figure 9. XRD patterns of (a) pristine P3-Na3/4Co1/3Ni1/3Mn1/3O2 and (b) oxide electrodes after the first charging to 3.5 V, (c) after the first charging to 3.5 V followed by discharging to 1.0 V, and (d) after 10 cycles between 1.0 and 3.5 V (stopped at 1.0 V). All P3Na3/4Co1/3Ni1/3Mn1/3O2 electrodes are tested in a full sodium−lithium cell. The Miller indexes for the P3- and O3-type of structure are indicated.

diffraction peaks after the electrochemical reaction, we examine the changes in the basal plane only (Figure 9). After the first charging up to 3.5 V, there is an increase in the interlayer space from 5.56 to 6.80 Å. The observed increase in the interlayer space of P3-Na3/4Co1/3Ni1/3Mn1/3O2 can be interpreted by taking into account the competition between the sodium content and electrostatic repulsion between charged Co1/3Ni1/3Mn1/3O2-layers: The smaller the sodium content, the higher the electrostatic repulsion. The interlayer expansion gives evidence for the ext raction of Na + from Na3/4Co1/3Ni1/3Mn1/3O2 during the charging of the full cell up to 3.5 V. During the reverse process of discharging of the full cell down to 1.0 V, the interlayer space decreases reaching a value of 5.59 Å, which is slightly higher than that of the pristine composition (Figure 9). Taking into account the correlation between the interlayer space and the sodium content,17 it appears that the oxide NaxCo1/3Ni1/3Mn1/3O2 with x = 0.5 is formed during the cell discharging. In addition, a broad peak at around 4.81 Å starts to develop. On the one hand, the strongly contracted interlayer space is related to intercalation of smaller Li + instead of bigger Na+ ions. On the other hand, the observation of phases with different interlayer spaces proves the intercalation of both Na + and Li+ into NaxCo1/3Ni1/3Mn1/3O2 during the full cell discharging. The two phases coexist even in the oxide cycled 10 times between 1.0 and 3.5 V, with the intensity of the phase with d =

Figure 10. Bright-field images/polycrystalline diffraction (top), HRTEM (middle), and bright field images/SAED (bottom) of cycled oxide P3-NaxCo1/3Ni1/3Mn1/3O2 after 10 cycles between 1.0 and 3.5 V (the cell test is stopped at 1.0 V, i.e., the oxide is in a charged state).

shows bright-field images and the corresponding polycrystalline electron diffraction for the oxide cycled 10 times between 1.0 and 3.5 V. As can be seen, the diffraction pattern consists of mixture of two layered phases. The first phase can be indexed in the R3̅m space group with lattice parameters of a = 0.284 nm and c = 1.44 nm and an interlayer space of 0.48 nm. This phase is due to a lithium-containing phase with three-layered stacking (O3-type of structure). For the sake of comparison, the lithium analogue LiCo1/3Ni1/3Mn1/3O2, obtained by a solid state reaction at 950 °C, is characterized by the same O3-type of structure but with slightly different lattice parameters: a = 2.8541 Å and c = 14.2130 Å.35 This proves once again the transformation of the pristine P3- into O3-phase during the lithium intercalation. The second phase has also layered 17329


Research Article


electrochemical reaction. However, the lower sodium content implies indirectly that the Li+ intercalation is the dominant reaction. In addition, the consistency in the ratio between Co, Ni, and Mn ions and their homogeneous distribution over particles provide evidence for the stability of transition metal layers during alkaline intercalation process. The participation of Na3/4Co1/3Ni1/3Mn1/3O2 in the electrochemical reaction is monitored by means of EPR. The EPR spectrum of pristine Na3/4Co1/3Ni1/3Mn1/3O2 consists of single Lorentzian line with a g-factor of 1.9862 and a line width of 159.4 mT (Figure 12). On the basis of the EPR study of

structure, but it displays higher interlayer space. This is consistent with XRD observation of the Na-containing phase with d = 0.56 nm (Figure 9). The most important finding is revealed by HRTEM analysis (Figure 10). The HRTEM images display lattice fringes with variable lattice spacing, where at least three different lattice spaces can be distinguished: 0.48, 0.52, and 0.88 nm. Although the first two lattice fringes are related to the interlayer spaces of the Li- and Na-containing phases, the latter lattice fringe is significantly higher (with more than 50%), and it cannot be explained on the basis of the electrostatic repulsion between charged transition metal layers. A large expansion of the interlayer space has been reported only when H2O molecules are inserted into the interlayer space of sodium-deficient transition metal oxides leading to the formation of hydrated phases with a general composition of NaxMO2.nH2O.36−38 The structure of hydrated oxides consists of mono- and bilayered Nax(H2O)y network separating the 2D MO2 layers by two distinct different distances of approximately 0.7 and 1.0 nm, respectively.36−38 The variation observed by us in the lattice spacing of the cycled oxide can be associated with a possible intercalation of organic molecules coming from the electrolyte solution: EC and/or DMC. However, it is worth taking into account that the simultaneous intercalation of bigger Na+ and smaller Li+ ions between basal planes will create planar defects related with variation of the interlayer spaces. Moreover, the appearance of planar defects is also in agreement with SAED taken along [010] direction, where oval diffraction spots are clearly observed (Figure 10). The composition of the cycled oxide P3NaxCo1/3Ni1/3Mn1/3O2 is evaluated by BF-STEM (Figure 11).

Figure 12. EPR spectra, temperature dependence of the g-factor and EPR line width (ΔHpp) for electrodes P3-Na3/4Co1/3Ni1/3Mn1/3O2 (a) after the first charging to 3.5 V and (b) after the first charging to 3.5 V followed by discharging to 1.0 V. Green lines correspond to the pristine composition.

Na x Co 1 / 3 Ni 1 / 3 Mn 1 / 3 O 2 (0.38 ≤ x ≤ 0.75) 1 7 and NaxNi1/2Mn1/2O2 (x = 0.50 and 0.67),19 the EPR response originated from Mn4+ ions, whereas the paramagnetic Ni2+ and Ni3+ ions have an impact on the temperature dependence of both the g-factor and the EPR line width. Upon cooling down from 300 to 100 K, there is a decrease in the g-factor and the EPR line width (Figure 11b). The EPR intensity increases with lowering of the registration temperature following the Curie− Weiss law with a Weiss constant of Θ = +(11 ± 7) K. After the Na extraction from P3-Na3/4Co1/3Ni1/3Mn1/3O2, the EPR signal undergoes a significant change with respect to the g-factor, line width, and temperature dependence of the signal intensity (Figure 12). The g-factor of desodiated oxide is lower than that of the pristine oxide in the whole registration temperature. The EPR line of desodiated oxide increases with decreasing the registration temperature; this behavior is opposite to that observed with the pristine oxide. In the temperature range of 100−300 K, the signal intensity obeys the Curie−Weiss law with a Weiss constant of Θ = −(60 ± 7) K. It is noticeable that, after oxide desodiation, the sign of the Weiss constant is changed from positive into negative. On the basis of EPR parameters, the EPR signal can be assigned to antiferromagnetically coupled Mn4+. The observed changes in the EPR signal of Mn4+ during the desodation process are associated with a corresponding variation in the oxidation state of transition metal ions, which are surrounding manganese ions. According to the charging/discharging curves, the extraction of Na+ is taking place due to the oxidation of both nickel and

Figure 11. BF-STEM images and corresponding composition map of CoKα1, Ni Kα1, Mn Kα1, O Kα1, and Na Kα1−2 for cycled electrodes NaxCo1/3Ni1/3Mn1/3O2.

After the electrochemical reaction, Co, Ni, and Mn ion contents remain unchanged: 0.35 mol for Co, 0.35 mol for Ni, and 0.30 mol for Mn versus 0.34 mol for Co, 0.31 mol for Ni, and 0.35 mol for Mn for the pristine oxide (not shown). It is noticeable that the Co, Ni, Mn, and O elements are homogeneously distributed over nanosized particles, subject to the electrochemical charge/discharge (Figure 11). Contrary to transition metal ions, the Na content sharply decreases and varies between 0.05 and 0.10 mol. This reveals unambiguously that Na+ ions are deintercalated and intercalated during the 17330


Research Article


1071.5 eV. The binding energy of the broad signal is shifted with more than 2 eV in comparison with that due to Na+ ions inserted in the layered structure (1070.2 eV, not shown). This allows us to assign the broad Na 1s signal to NaF, which can originate either from the interaction of Na3/4Co1/3Ni1/3Mn1/3O2 with the PVDF binder during the electrode manufacturing (as it was observed with Li4Ti5O12) or from electrolyte−electrode interaction. Contrary to the Na 1s spectrum, at least two overlapping signals can be resolved in the Na 2s energy region (Figure 13). The signal with a binding energy of 63.9 eV complementary to the Na 1s signal at 1071.5 eV is due to NaF.27 The second signal at 61.4 eV comes from Na+ ions belonging to layered oxide. The observation of the Na 2s signal and the lack of the Na 1s one is related to different analysis depths. This means that few top surface layers of oxide NaxCo1/3Ni1/3Mn1/3O2 are covered smoothly by NaF. It is worth mentioning that this is not observed with the cycled electrode Li4Ti5O12, where both NaF and Na-containing phase coexist simultaneously even in the few top layers. In accordance with Na 1s and Na 2s, the P 2p spectrum consists of broad signal with binding energy of 137.3 eV due to NaxPFy compounds (Figure 13). In the F 1s spectrum, the signal coming from fluoride compounds is also observed, i.e., binding energy of 685.1 eV. The formation of Na xPFy compounds is an indication of interaction between the oxide surface and lithium electrolyte. In addition, fluorine atoms in PVDF binder are responsible for the appearance of the intensive F 1s signal with binding energy of 687.2 eV. The most important result is the lack of any signal due to Li, which confirms once again that during the charging of the full sell up to 3.5 V only sodium is being extracted from Na3/4Co1/3Ni1/3Mn1/3O2 (Figure 13). The extracted sodium passes into the lithium electrolyte, where it demonstrates reactivity toward Li4Ti5O12. During the reverse process of discharging the full cell down to 1.0 V, the Na 1s, Na 2s, P 2p, and F 1s spectra do not show any measurable differences (Figure 13). In contrast, two new signals at 56.8 and 54.4 eV are growing in intensities in the Li 1s spectrum. The signal at 56.8 eV corresponds to Li+ in LiF, whereas the signal at 54.4 eV is a result from Li+ ions bonded to oxygen in the layered phase. This result means that the discharging of the full cell initiates the interaction of oxide surface with lithium electrolyte, as well as the intercalation of Li+ into layered oxide. The prolonged cycling leads to an increase in the relative amount of NaF versus Na-containing phase. It is important that the layered oxide is more reactive toward the lithium electrolyte in comparison with Li4Ti5O12 negative electrode.

cobalt ions. Although the oxidation of paramagnetic nickel ions, i.e., Ni2+ and Ni3+, produces diamagnetic Ni4+ ions (S = 0), the diamagnetic Co3+ ions are oxidized into diamagnetic Co4+ ones. The changes in the spin states of Ni and Co ions are reflected in the temperature dependence of both the g-factor and the EPR line width (Figure 12). The reverse process of intercalation of alkaline ions into the layered oxide does not restore the EPR signal registered with the pristine oxide (Figure 12). At 300 K, the g-factor of intercalated oxide is close to that of the desodiated oxide, but it displays a slighter dependence on the registration temperature. The EPR line decreases slightly after alkaline ion intercalation with preservation of the trend to line broadening with decreasing the registration temperature. The magnitude of the Weiss constant becomes smaller, but its sign remains negative: Θ = −(21 ± 9) K. This means it is the Mn4+ ions that are responsible for the appearance of the EPR signal come from cycled layered oxide. However, the observation of Mn4+ ions in discharged oxide shows their electrochemical inactivity during the electrochemical process. This is in a good accordance with charging/discharging experiments (Figure 3). Surface Changes in P3-Na3/4Co1/3Ni1/3Mn1/3O2 during Charging/Discharging Process. To identify the surface composition of P3-Na3/4Co1/3Ni1/3Mn1/3O2, Figure 13 compares the XPS spectra of the three types of cycled electrodes in the Na 1s, Na 2s, Li 1s, F 1s, and P 2p regions. For the desodiated oxide, i.e., for the oxide charged up to 3.5 V, the Na 1s spectrum displays a broad signal with a center of gravity at

■

CONCLUSIONS We demonstrate a manufacturing of new type of hybrid sodium−lithium ion cell by using a unique electrode combination (Li 4 T i 5 O 1 2 spinel and layered P3Na3/4Co1/3Ni1/3Mn1/3O2 as a negative and positive electrodes) and conventional lithium electrolyte (LiPF6 salt dissolved in EC/DMC). Both Li4Ti5O12 and P3-Na3/4Co1/3Ni1/3Mn1/3O2 electrodes are fabricated by using PVDF as a binder, which is involved in an interaction with surface oxides. The cell operates at an average voltage of about 2.35 V by delivering about 100 mAh/g. The electrochemical process is based on the Ni2+/ Ni3+,4+, Co3+/Co4+, and Ti3+/Ti4+ redox pairs, and it includes a competitive intercalation of Na+ and Li+ into negative and

Figure 13. XPS spectra in the Na 1s, Na 2s, Li 1s, F 1s and P 2p regions for (a) P3-Na3/4Co1/3Ni1/3Mn1/3O2 positive electrodes after the first charging of the full cell to 3.5 V, (b) after the first cycle of the full cell to 1.0 V, (c) and after 10 cycles between 1.0−3.5 V (the cell is stopped at 1.0 V). 17331



■

positive electrode oxides. During the whole process, Mn4+ ions are remaining electrochemically inactive. The electrochemical behavior of the sodium−lithium full cell is different from those for the half-cells versus Li metal anode. The cell function starts with a charging mode, during which 1/3 mol of Na+ ions are being extracted from layered P3Na3/4Co1/3Ni1/3Mn1/3O2 upon preserving the layered structure matrix. The extracted Na+ ions together with Li+ ions from electrolyte are being inserted into Li4Ti5O12 spinel leading to its transformation into cubic Li7Ti5O12 and Li7−xNaxTi5O12 (where x tends to 1 mol),; the total amount of inserted ions was found to be 2/3 mol. The ratio of Na-intercalated cubic phase is higher than that in the Li one. The bulk phase changes are concomitant with corresponding surface alteration: The few top surface layers of P3-Na3/4Co1/3Ni1/3Mn1/3O2 are covered smoothly with NaxPyFz, whereas both LixPyFz and NaxPyFz are unhomogeneously deposed on the surface of Li4Ti5O12. After the reverse process of discharging, there is a simultaneous extraction of Li+ and Na+ ions occurring from the negative electrode leading to a restoration of Li4Ti5O12 spinel phase. However, some sodium-containing cubic phase coexists with Li4Ti5O12. The released Na+ and Li+ ions are simultaneously being inserted into layered NaxCo1/3Ni1/3Mn1/3O2 oxide, as a result of which sodium-and lithium-rich phases are formed. The discharging of the full cell up to 1 V intensifies the interaction of the surface of both layered and spinel phases with lithium electrolyte. On prolonged cycling, the intercalation of Li+ becomes the dominant process over Na+ intercalation, which is valid for both the layered and spinel phase. The simultaneous Li+ and Na+ intercalations lead to the occurrence of planar defects in the layered oxides, whereas in the case of the spinel phase, there appear closely interconnected nanodomains with different lattice parameters. The layered oxide is more reactive toward the lithium electrolyte in comparison with the Li4Ti5O12 spinel. To the best of our knowledge, the present work is the first report on the fabrication of a full hybrid sodium−lithium ion cell, which operates at 2.35 V by using a conventional organic lithium electrolyte. This type of battery differs from the proposed earlier mixed Li+/Na+ batteries with positive and negative electrodes such as LiMn2O 4/Na0.22MnO 2 and Na0.44MnO2/TiP2O7, respectively, and an aqueous electrolyte that operates between 0.2 and 1.05 V.23 The difference comes from the reactivity of the selected electrodes in respect of Li+ and Na+ ions: While both Li4Ti5O12 spinel and layered P3Na3/4Co1/3Ni1/3Mn1/3O2 oxide can intercalate Li+ and Na+ simultaneously, LiMn2O4 spinel and Na0.22MnO2 oxides react preferentially only with Li+ and only with Na+, respectively. Although the electrochemical performance of the sodium− lithium full cell proposed by us is not yet optimized, these first studies could initiate further investigations focused on the design of full cells involving lithium and sodium compounds as negative and positive electrodes.

■

Research Article

AUTHOR INFORMATION

Corresponding Author

*Phone: +359 2 9793915. Fax: +359 2 8705024. E-mail: [email protected]. Notes

The authors declare no competing financial interest.

■

REFERENCES

(1) Thackeray, M. M. Structural Considerations of Layered and Spinel Lithiated Oxides for Lithium Ion Batteries. J. Electrochem. Soc. 1995, 142, 2558−2563. (2) Nagaura, T.; Tozawa, K. Lithium Ion Rechargeable Battery. Prog. Batteries Solar Cells 1990, 9, 209−217. (3) Dunn, B.; Kamath, H.; Tarascon, J. M. Electrical Energy Storage for the Grid: A Battery of Choices. Science 2011, 334, 928−935. (4) Yi, T.-F.; Yang, S.-Y.; Xie, Y. Recent Advances of Li4Ti5O12 as Promising Next Generation Anode Material for High Power LithiumIon Batteries. J. Mater. Chem. A 2015, 3, 5750−5777. (5) Hong, J.; Gwon, H.; Jung, S.-K.; Ku, K.; Kang, K. Review Lithium-Excess Layered Cathodes for Lithium Rechargeable Batteries. J. Electrochem. Soc. 2015, 162, A2447−A2467. (6) Ohzuku, T.; Ueda, A.; Yamamoto, N. Zero-Strain Insertion Material of Li[Li1/3Ti5/3]O4 for Rechargeable Lithium Cells. J. Electrochem. Soc. 1995, 142, 1431−1435. (7) Lu, W.; Liu, J.; Sun, Y.-K.; Amine, K. Electrochemical Performance of Li4/3Ti5/3O4/Li1+x(Ni1/3Co1/3Mn1/3)1−xO2 Cell for High Power Applications. J. Power Sources 2007, 167, 212−216. (8) Guo, X.; Xiang, H. F.; Zhou, T. P.; Li, W. H.; Wang, X. W.; Zhou, J. X.; Yu, Y. Solid-State Synthesis and Electrochemical Performance of Li4Ti5O12/graphene Composite for Lithium-Ion Batteries. Electrochim. Acta 2013, 109, 33−38. (9) Ellis, B. L.; Nazar, L. F. Sodium and Sodium-Ion Energy Storage Batteries. Curr. Opin. Solid State Mater. Sci. 2012, 16, 168−177. (10) Yabuuchi, N.; Kubota, K.; Dahbi, M.; Komaba, S. Research Development on Sodium-Ion Batteries. Chem. Rev. 2014, 114, 11636− 11682. (11) Han, M. H.; Gonzalo, E.; Singh, G.; Rojo, T. A Comprehensive Review of Sodium Layered Oxides: Powerful Cathodes for Na-Ion Batteries. Energy Environ. Sci. 2015, 8, 81−102. (12) Delmas, C.; Fouassier, C.; Hagenmuller, P. Structural Classification and Properties of the Layered Oxides. Physica B+C 1980, 99, 81−85. (13) Senguttuvan, P.; Rousse, G.; Seznec, V.; Tarascon, J.-M.; Palacín, M. R. Na 2 Ti 3 O 7: Lowest Voltage Ever Reported Oxide Insertion Electrode for Sodium Ion Batteries. Chem. Mater. 2011, 23, 4109−4111. (14) Xu, J.; Ma, C.; Balasubramanian, M.; Meng, Y.-S. Understanding Na2Ti3O7 as an Ultra-Low Voltage Anode Material for a Na-Ion Battery. Chem. Commun. 2014, 50, 12564−12567. (15) Kalapsazova, M.; Stoyanova, R.; Zhecheva, E. Structural Characterization and Electrochemical Intercalation of Li+ in Layered Na0.65Ni0.5Mn0.5O2 Obtained by Freeze-Drying Method. J. Solid State Electrochem. 2014, 18, 2343−2350. (16) Yoncheva, M.; Stoyanova, R.; Zhecheva, E.; Kuzmanova, E.; Sendova-Vassileva, M.; Nihtianova, D.; Carlier, D.; Guignard, M.; Delmas, C. Structure and Reversible Lithium Intercalation in a New P’3-phase: Na2/3Mn1−yFeyO2 (y = 0, 1/3, 2/3). J. Mater. Chem. 2012, 22, 23418−23427. (17) Ivanova, Sv.; Zhecheva, E.; Kukeva, R.; Tyuliev, G.; Nihtianova, D.; Mihailov, L.; Stoyanova, R. Effect of Sodium Content on the Reversible Lithium Intercalation into Sodium-deficient Cobalt-NickelManganese Oxides NaxCo1/3Ni1/3Mn1/3O2 (0.38< x ≤.75) with a P3type of Structure. J. Phys. Chem. C 2016, 120, 3654−3668. (18) Kalapsazova, M.; Ortiz, G. F.; Tirado, J. L.; Dolotko, O.; Zhecheva, E.; Nihtianova, D.; Mihaylov, L.; Stoyanova, R. P3-type Layered Sodium-Deficient Nickel-Manganese Oxides: a Flexible Structural Matrix for Reversible Sodium and Lithium Intercalation. ChemPlusChem 2015, 80, 1642−1656.

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.6b05075. XRD patterns of Li 4 Ti 5 O 1 2 spinel and P3− Na0.75Co1/3Ni1/3Mn1/3O2 (PDF) 17332


Research Article


conducting Sodium Cobalt Oxide and its Related Phase. J. Solid State Chem. 2004, 177, 372−376. (38) Park, S.; Lee, Y.; Elcombe, M.; Vogt, T. Synthesis and Structure of the Bilayer Hydrate Na0.3NiO2·1.3D2O. Inorg. Chem. 2006, 45, 3490−3492.

(19) Kalapsazova, M.; Stoyanova, R.; Zhecheva, E.; Tyuliev, G.; Nihtianova, D. Sodium Deficient Nickel-Manganese Oxides as Intercalation Electrodes in Lithium Ion Batteries. J. Mater. Chem. A 2014, 2, 19383−19395. (20) Sun, Y.; Zhao, L.; Pan, H.; Lu, X.; Gu, L.; Hu, Y.-S.; Li, H.; Armand, M.; Ikuhara, Y.; Chen, L.; Huang, X. Direct Atomic-scale Confirmation of Three-Phase Storage Mechanism in Li4Ti5O12 Anodes for Room-Temperature Sodium-Ion Batteries. Nat. Commun. 2013, 4, 1870. (21) Yu, X.; Pan, H.; Wan, W.; Ma, C.; Bai, J.; Meng, Q.; Ehrlich, S. N.; Hu, Y. S.; Yang, X. Q. A Size-Dependent Sodium Storage Mechanism in Li4Ti5O12 Investigated by a Novel Characterization Technique Combining in situ X-ray Diffraction and Chemical Sodiation. Nano Lett. 2013, 13, 4721−4727. (22) Yao, H.-R.; You, Y.; Yin, Y.-X.; Wan, L.-J.; Guo, Y.-G. Rechargeable Dual-Metal-Ion Batteries for Advanced Energy Storage. Phys. Chem. Chem. Phys. 2016, 18, 9326−9333. (23) Chen, L.; Gu, Q.; Zhou, X.; Lee, S.; Xia, Y.; Liu, Z. New-concept Batteries Based on Aqueous Li+/Na+ Mixed-ion Electrolytes. Sci. Rep. 2013, 3, 1946. (24) Rodrıguez-Carvajal, J. Commission on Powder Diffraction. Newsletter 2001, 26, 12−19. (25) Kitta, M.; Kuratani, K.; Tabuchi, M.; Takeichi, N.; Akita, T.; Kiyobayashi, T.; Kohyama, M. Irreversible Structural Change of a Spinel Li4Ti5O12 Particle via Na Insertion-Extraction Cycles of a Sodium-Ion Battery. Electrochim. Acta 2014, 148, 175−179. (26) Song, M.-S.; Kim, R.-H.; Baek, S.-W.; Lee, K.-S.; Park, K.; Benayad, A. Is Li4Ti5O12 a Solid-Electrolyte-Interphase-Free Electrode Material in Li-Ion Batteries? Reactivity between the Li4Ti5O12 Electrode and Electrolyte. J. Mater. Chem. A 2014, 2, 631−636. (27) Ismail, I.; Noda, A.; Nishimoto, A.; Watanabe, M. XPS Study of Lithium Surface after Contact with Lithium-Salt Doped Polymer Electrolytes. Electrochim. Acta 2001, 46, 1595−1603. (28) Erdem, B.; Hunsicker, R. A.; Simmons, G. W.; Sudol, E. D.; Dimonie, V. L.; El-Aasser, M. S. XPS and FTIR Surface Characterization of TiO2 Particles Used in Polymer Encapsulation. Langmuir 2001, 17, 2664−2669. (29) Lee, B.-R.; Oh, E.-O. Effect of Molecular Weight and Degree of Substitution of a Sodium-Carboxymethyl Cellulose Binder on Li4Ti5O12 Anodic Performance. J. Phys. Chem. C 2013, 117, 4404− 4409. (30) Brisson, P.-Y.; Darmstadt, H.; Fafard, M.; Adnot, A.; Servant, G.; Soucy, G. X-ray Photoelectron Spectroscopy Study of Sodium Reactions in Carbon Cathode Blocks of Aluminium Oxide Reduction Cells. Carbon 2006, 44, 1438−1447. (31) Andersson, A. M.; Herstedt, M.; Bishop, A.; Edström, K. The Influence of Lithium Salt on the Interfacial Reactions Controlling the Thermal Stability of Graphite Anodes. Electrochim. Acta 2002, 47, 1885−1898. (32) Eriksson, T.; Andersson, A. M.; Gejke, G.; Gustafsson, T.; Thomas, J. O. Influence of Temperature on the Interface Chemistry of LixMn2O4 Electrodes. Langmuir 2002, 18, 3609−3619. (33) Edström, K.; Gustafsson, T.; Thomas, J. O. The Cathode− Electrolyte Interface in the Li-Ion Battery. Electrochim. Acta 2004, 50, 397−403. (34) Dedryvère, R.; Leroy, S.; Martinez, H.; Blanchard, F.; Lemordant, D.; Gonbeau, D. XPS Valence Characterization of Lithium Salts as a Tool to Study Electrode/Electrolyte Interfaces of Li-Ion Batteries. J. Phys. Chem. B 2006, 110, 12986−12992. (35) Shinova, E.; Stoyanova, R.; Zhecheva, E.; Ortiz, G. F.; Lavela, P.; Tirado, J. L. Cationic Distribution and Electrochemical Performance of LiCo1/3Ni1/3Mn1/3O2 Electrodes for Lithium-Ion Batteries. Solid State Ionics 2008, 179, 2198−2208. (36) Takada, K.; Sakurai, H.; Takayama-Muromachi, E.; Izumi, F.; Dilanian, R. A.; Sasaki, T. Superconductivity in Two-Dimensional CoO2 Layers. Nature 2003, 422, 53−55. (37) Takada, K.; Sakurai, H.; Takayama-Muromachi, E.; Izumi, F.; Dilanian, R. A.; Sasaki, T. Structural Difference between a Super17333


3O2 versus ... - ACS Publications

Recommend Documents