3Mn0.8Fe0.1Ti0.1O2 Cathode

Article pubs.acs.org/cm

High-Performance P2-Phase Na2/3Mn0.8Fe0.1Ti0.1O2 Cathode Material for Ambient-Temperature Sodium-Ion Batteries Man Huon Han,*,† Elena Gonzalo,† Neeraj Sharma,§ Juan Miguel López del Amo,† Michel Armand,† Maxim Avdeev,∥ Jose Javier Saiz Garitaonandia,‡ and Teófilo Rojo*,†,‡ CIC EnergiGUNE, Parque Tecnológico de Á lava, Albert Einstein 48, Edificio CIC, 01510 Miñano, Á lava, Spain Facultad de Ciencia y Tecnología, Universidad del País Vasco/Euskal Herriko Unibertsitatea, Post Office Box 644, 48080 Bilbao, Spain § School of Chemistry, University of New South Wales, Sydney, New South Wales 2052, Australia ∥ Australian Nuclear and Science Technology Organisation, Locked Bag 2001, Kirrawee DC, New South Wales 2232, Australia † ‡

S Supporting Information *

ABSTRACT: High-performance Mn-rich P2-phase Na2/3Mn0.8Fe0.1Ti0.1O2 is synthesized by a ceramic method, and its stable electrochemical performance is demonstrated. 23Na solid-state NMR confirms the substitution of Ti4+ ions in the transition metal oxide layer and very fast Na+ mobility in the interlayer space. The pristine electrode delivers a second charge/discharge capacity of 146.57/144.16 mA·h·g−1 and retains 95.09% of discharge capacity at the 50th cycle within the voltage range 4.0−2.0 V at C/10. At 1C, the reversible specific capacity still reaches 99.40 mA·h·g−1, and capacity retention of 87.70% is achieved from second to 300th cycle. In addition, the moisture-exposed electrode reaches reversible capacities of more than 130 and 80 mA·h·g−1 for C/10 and 1C, respectively, with excellent capacity retention. The correlation between overall electrochemical performance of both electrodes and crystal structural characteristics are investigated by neutron powder diffraction. The stability of pristine electrode’s crystallographic structure during the charge/discharge process has been investigated by in situ Xray diffraction, where only a solid solution reaction occurs within the given voltage range except for a small biphasic mechanism occurring at or below 2.2 V during the discharge process. The relatively small substitution (20%) at the transition metal site leads to stable electrochemical performance, which is in part derived from the structural stability during electrochemical cycling. Therefore, the small cosubstitution (e.g., with Ti and Fe) route suggests a possible new scope for the design of sodium-ion battery electrodes that are suitable for long-term cycling.

1. INTRODUCTION Commercialization of the Li-ion battery (LIB) during the last decades has been a major success in the secondary battery industry and has made it possible to implement small electronic gadgets into daily life. Today, the application of LIB is being explored for electric vehicles (EV) and beyond.1 However, expanding the use of Li ion-based batteries to larger applications such as electricity storage grid support may be hindered by the availability of Li sources.2 In this regard, Na-ion batteries (NIB) are considered as an excellent alternative to LIB for stationary applications because sodium is one of the most abundant elements in Earth’s crust and seawater3 and is also the second lightest alkali metal next to Li. In fact, the Na intercalation chemistry of some cathode materials for NIB has been demonstrated during the 1980s,4−7 and recent research progress in electrode materials has shown the feasibility of costeffective fully functional units that can be implemented into real applications in the near future.8−12 It is widely postulated that the energy density of NIB cathode materials may not be able to match that of LIB, because Na is more than 3 times heavier © 2015 American Chemical Society

than Li and the average operating voltage of intercalation materials is lower than that of Li-host counterparts by 300−500 mV.13 Despite such drawbacks, NIBs are an attractive system because the cost and durability for long-term operation are more important than the energy density of each unit for large stationary applications.14 Furthermore, in NIB systems, cheaper aluminum current collectors can be used at both electrodes (≠Cu), a considerable advantage in terms of sustainability. Among the cathode candidates, layered oxides NaxMTO2 (MT = Ti, V, Cr, Mn, Fe, Co, Ni, and combinations of two or more MT) are of great interest because of their large specific capacity, ease of synthesis, and choice of various metal constituents.15 In particular, the series of Fe- and Mncontaining layered oxides NaxFe1−yMnyO2 (x ≤ 1; 0 ≤ y ≤ 1) have been investigated extensively16−19 due to their environmental benignity, inexpensive metal constituents, and Received: August 24, 2015 Revised: November 24, 2015 Published: November 25, 2015 106

DOI: 10.1021/acs.chemmater.5b03276 Chem. Mater. 2016, 28, 106−116

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Chemistry of Materials

Figure 1. Rietveld analysis of pristine P2-phase Na2/3Mn0.8Fe0.1Ti0.1O2 powder by use of FullProf software package with hexagonal P63/mmc symmetry. (Inset) Structural illustration of a typical P2-phase layered oxide.

large reversible capacity. α-NaFeO2 exhibits a high, flat voltage profile during electrochemical cycling but undergoes an irreversible structural change above approximately 3.6 V, which causes rapid capacity degradation after only a few cycles.20 Obvious downsides of Mn-containing layered oxides are the possibility of Mn ion dissolution in electrolyte21 and rapid capacity fade due to structural strain caused by introduction of Jahn−Teller high-spin Mn3+ (t2g3eg1) during electrochemical cycling despite the large initial reversible capacity and low polarization. 7,22,23 Recently, a small substitution of an electrochemically inactive element into the Mn-rich P2 phase has shown stable cycling performance.24,25 Although the low voltage profile of the Mn redox potential (low energy density) still remains an obstacle, substitution of electrochemically inactive elements into the oxide layer is a viable approach to achieve stable electrochemical performance for Mn-rich layered oxides. For the loss of a small amount of energy density (capacity or active redox couple), the electrode can be stabilized to cycle for longer periods of time. In the present study, the electrochemical performance and structural stability of partially substituted P2-phase Na2/3Mn0.8Fe0.1Ti0.1O2 are investigated. As evidenced by a peak shift, the 23Na NMR spectrum confirms the substitution of Ti4+ into the oxide layer. The pristine electrode delivers a second charge/discharge capacity of 146.57/144.16 mA·h·g−1 within the voltage range 4.0−2.0 V at C/10. The capacity retention reaches 95.09% after 50 cycles at C/10 in Na half-cell configuration, and an excellent rate capability is achieved with 99.40 mA·h·g−1 discharge capacity at 1C (68.95% of specific capacity of C/10). In addition, for cycling at 1C, the capacity retention reaches 87.70% after 300 cycles. Such high capacity, stable capacity retention, and high rate capability are comparable to or better than those of high-performance layered oxide electrodes published to date.12 In addition, the reversible capacity of a moisture-exposed version of this electrode reaches 130 and 80 mA·h·g−1 for C/10 and 1C, respectively, with similar or better capacity retention after the few first cycles. Such superior electrochemical performance of the pristine and moisture-exposed Na2/3Mn0.8Fe0.1Ti0.1O2 electrodes without optimizing synthesis conditions, electrode preparation, and cell components are indicative of outstanding stability of the present electrode within the experimental

conditions. Detailed structural analysis of the pristine and moisture-exposed phase is performed with neutron powder diffraction data and the structural stability of pristine electrode is demonstrated by in situ X-ray diffraction (XRD), where only a solid solution region is observed during the charge process and a small biphasic region is found at the end of the discharge process.

2. EXPERIMENTAL SECTION P2-phase Na2/3Mn0.8Fe0.1Ti0.1O2 was synthesized by the ceramic method. A stoichiometric mixture of Na2CO3, Mn2O3, Fe2O3, and TiO2 powders was mixed by use of a speed mixer in aqueous medium for 20 min at 1000 rpm. The mixture was quickly frozen by pouring liquid nitrogen directly into the mixing container and then dried in the freeze-dryer for 2 days. Pellets of pressed freeze-dried powders (diameter 20 mm, thickness 3 mm) were annealed at 1000 °C for 6 h under ambient atmosphere, and upon furnace cooling, the pellet was quickly transferred to the glovebox in order to avoid moisture contact. For the moisture-exposed sample, 100 mg of pristine powder was dispersed in 10 mL of deionized (DI) water for 24 h. The powder was filtered and washed with 100 mL of DI water. Then the powder was immediately frozen with liquid nitrogen and dried in a freeze-dryer for 1 day. The dried powder was transferred to an Ar-filled glovebox. For the sample for neutron powder diffraction studies, D2O was used instead of water. X-ray diffraction patterns of the powder samples were collected on a Bruker D8 Advance instrument with Cu Kα radiation of λ = 1.5418 Å within the 2θ range 15−80°. All samples were mounted onto an atmosphere-protective XRD sample holder with a Kapton film cover. In situ XRD cell was prepared from an in-house-designed Swagelok cell with a Be window and poly(oxymethylene) body. The cell was galvanostatically cycled over the voltage range 4.0−2.0 V at C/50 by use of a Biologic SP200 potentiostat. Each scan was collected in 0.02° increments between 15° and 50° at a scan speed of 0.02°/s. Elemental analysis was performed by inductively coupled plasma optical emission spectrometry (ICP-OES) on a Horiba Scientific Ultima 2 spectrometer and scanning electron microscopy (SEM) on a FEI Quanta 250 microscope equipped with an energy-dispersive X-ray spectroscopic detector (EDS). The elemental ratio calculation is based on the 0.8 mol Mn concentration. 23 Na magic-angle spinning nuclear magnetic resonance (MAS NMR) experiments were performed at 52.9 MHz on a Bruker-300 spectrometer charged to a field of 4.7 T using a 1.3 mm MAS probe at rotor spinning speed of 50 kHz. A rotor synchronized spin−echo pulse sequence (90°−τ−180°−τ1−acquisition) was used with typical 90° 107


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Figure 2. (a) Selected 2θ region of as-collected neutron powder diffraction patterns of pristine and D2O (moisture-exposed) samples. (b) Rietveld refined fit of pristine Na0.64(2)Mn0.8Fe0.1Ti0.1O2 (top) and D2O-treated Na0.52(2)Mn0.8Fe0.1Ti0.1O2 (bottom) models to neutron powder diffraction data, with data shown as crosses, calculated Rietveld model as a line through the data, difference between data and model as a line below the data, and vertical reflection markers for the phases modeled. (c) Crystal structure of D2O-treated sample. Mn/Ti/Fe are purple/blue/orange, oxygen is red, D site is labeled, and Na is yellow. The amount of shading on a sphere is related to the occupancy of the site. and 180° pulses of 1.2 and 2.4 μs, respectively, and recycle delay of 0.8 s. The spectra were referenced to a 0.1 M solution of 23NaCl. Neutron powder diffraction (NPD) data were obtained at the Open-Pool Australian Light-water (OPAL) Research Reactor, Australian Nuclear Science and Technology Organisation (ANSTO), by use of the high-resolution diffractometer Echidna.26 The samples were treated as air-sensitive powders and were sealed in 9 mm diameter vanadium cans. Room-temperature data were collected at a wavelength of λ = 1.6215(2) Å. Note the D2O (moisture-exposed) sample was run in a cryofurnace. Rietveld refinements were carried out by the GSAS27 software suite with EXPGUI28 software interface. Mössbauer spectroscopy was performed at room temperature in the transmission geometry on a conventional constant-acceleration spectrometer with a 57Co−Rh source. Laminate electrode preparation is carried out in an Ar-filled glovebox. The slurry electrode was prepared by mixing the active material, super carbon C65, and PVDF [poly(vinylidene difluoride)] in a mass ratio of 80:10:10 in NMP (N-methyl-2-pyrrolidone), followed by vigorous stirring for 2 h. The slurry was then cast onto an aluminum current collector sheet by use of a mini coater. The laminate was vacuum-dried at 120 °C for 12 h. Circular electrodes were punched out and pressed at 5 tons prior to battery assembly. Na halfcells were constructed with 1.0 M NaPF6 in a mixture of ethylene carbonate (EC)/propylene carbonate/fluoroethylene carbonate (FEC) in 49:49:2 ratio, and metallic Na was used as the anode in 2032 coin cells. Assembly was also undertaken in an Ar-filled glovebox. Galvanostatic electrochemical testing was performed with a BioLogic VMP3 within the voltage range 4.0−2.0 V. All the C-rates are calculated on the basis of full theoretical specific capacity of 245.19 mA·h·g−1 of 1 Na deintercalation/intercalation.

The average thickness and mass loading of dried laminate on 12 mm electrode was 30−35 μm and 1.2−1.4 mg, respectively.

3. RESULTS The P2-phase layered oxide is characterized by layered stacking of the oxide layer as AABBAA..., where Na+ occupies trigonal prismatic sites and shares either an entire edge or an entire face with adjacent oxide layers as seen in the inset of Figure 1. The XRD p attern matching of the pristine powder, Na2/3Mn0.8Fe0.1Ti0.1O2, in Figure 1 shows an excellent fit by use of hexagonal P63/mmc symmetry for the Bragg positions and observed peaks without any noticeable impurity peaks. The lattice parameters were calculated to be a = 2.9143(2) Å and c = 11.184(1) Å, which are similar to those of the unsubstituted NaxMnO2 compound [a = 2.8603(4) Å and c = 11.153(2) Å].21 A deviation in the a lattice parameter compared to the unsubstituted compound is due to introduction of the larger ionic radius Ti4+ (Mn3+ 0.645 Å, Mn4+ 0.53 Å, Fe3+ 0.645 Å, Ti4+ 0.605 Å; Mn and Fe are in the high-spin state with 6coordination)29 into the oxide layer. A small difference in the c lattice parameter reflects the slight difference in Na concentration in the pristine state, where a lower Na concentration results in larger interlayer spacing due to increased repulsion between adjacent oxide layers because the shielding effect by the positively charged Na ions decreases. Moisture exposure of the layered oxide is known to cause Na+/ H+ exchange to form a protonated phase,30−32 and further exposure to moisture promotes formation of the hydrated 108


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Chemistry of Materials phase, where H2O molecules occupy the interlayer space.33,34 The sample treated with deionized water will be referred to as moisture-exposed sample. The XRD pattern of the moistureexposed sample, refined by use of hexagonal P63/mmc symmetry with both the protonated and hydrated models assumed (see peak at 2θ = 25.3°), resulted in a good fit with the observed pattern without any impurity (see Figure S1 in Supporting Information). The c lattice parameter of the protonated model is slightly larger than the pristine phase, which is initially attributed to the smaller proton imposing a lower shielding effect than the larger Na+ ion, causing a stronger repulsion between adjacent oxide layers. The c lattice parameter of the hydrated model is significantly larger than the other phases because a large water molecule expands the interlayer space.29,35,36 A slight difference in the a lattice reflects changes in Mn4+/Mn3+ ratio in the transition metal layers upon Na+ leaching and/or protonation as discussed below. Moreover, recent literature suggests that deviation in the a lattice parameter can also be attributed to the insertion of CO32− in the interlayer space.33 To obtain details about the transition metal ratios and interlayer occupancy of Na and/or water or H, neutron powder diffraction analysis has been undertaken. The sample treated with deuterated water will be referred as D2O sample. The P2phase Na0.66MO2 structural model with P63/mmc symmetry was used as a starting point. A stepwise approach was then used to model the structure in each case. Differences in modeling the pristine and D2O samples with neutron powder diffraction data included changes in peak shapes and increased number of background terms required for the D2O sample. The latter is expected, as any of trace H2O in D2O significantly increases background due to larger incoherent neutron scattering cross section of H compared to D. Figure 2a shows a zoomed-in region of the as-collected pattern showing the slight change in the lattice (reflection 2θ value) between the samples, Figure 2b shows the Rietveld refined fits for pristine and D2O samples, Figure 2c shows the crystal structure model of the D2O sample, and Table S2 in Supporting Information lists the refined crystal structural information. The D2O sample lattice is expanded in the c or stacking axis direction by 0.4% (see Table S2 in Supporting Information), which suggests that changes caused by the D2O synthetic procedure are minimal in terms of lattice expansion. This also indicates that D2O insertion does not happen in the interlayer spacing, as the expected c lattice parameter for such a case would be closer to ∼14 Å as shown in NaxCoO2·y′ H2O.37 Furthermore, the XRD results suggest that the majority phase, protonated, has a c lattice parameter closer to that of pristine phase, which correlates with the main phase observed in neutron powder diffraction data and its c lattice parameter. This indicates some Na+/D+ exchange might be occurring. From the Lorentzian peak-shape profile term in combination with the Scherrer equation, the particle size distribution observed by neutron powder diffraction is much smaller for the D2O sample at 16 nm, compared to 70 nm for the pristine sample. Note the measured particle size distribution is correlated to the scattering planes (and particles), and the SEM data shown in Figure 3 (e.g., secondary particles) are more representative of the electrode surface. In terms of atomic parameters, a number of models were tried. The refined Na content in the D2O sample was found to be systematically lower than that in the pristine sample, 0.52(2) and 0.64(2), respectively, and this agrees with the Na content

Figure 3. SEM micrographs of (a) pristine P2-phase Na2/3Mn0.8Fe0.1Ti0.1O2 powder, (b) moisture-exposed powder, and (c) EDS mapping of pristine electrode. (a, inset) Particle used for EDS mapping.

derived from ICP and EDS analyses. On the basis of the model for brucite Mg(OD)2,38 D+ was positioned ∼0.96 Å below the oxygen site, forming a hydroxyl (OD) grouping. The refined D occupancy was 0.022(4), or around 4% D + in the Na0.52(2)D0.04(1)Mn0.8Fe0.1Ti0.1O2 compound. Irregularly shaped micrometric particles exhibit a relatively large size distribution for both pristine and moisture-exposed samples as seen in Figure 3a,b. EDS and ICP elemental analysis indicates that the Mn:Fe:Ti ratio is close to the theoretical value of 0.8:0.1:0.1, while the Na concentration varies upon moisture exposure (see Table S3 in Supporting Information). The homogeneity of each element throughout the particle is shown by EDS mapping on a single particle, shown in Figure 3c. Na concentration analyzed by ICP for pristine powder is slightly larger than EDS for pristine electrode, which may be due to Na leaching during sample handling by water contamination. However, ICP and EDS analysis of the Na concentration for moisture-exposed electrode yielded identical Na concentration of 0.43Na, which is close to the value of other P2-phase layered oxides that were exposed for significantly shorter times to moisture.29−31 The observation of lower Na concentration in the moisture-exposed sample suggests that Na leaching by neutral water is limited and stabilizes to approximately 0.43Na for the Mn-containing P2-phase material. Notably, the lower Na value derived for the moisture-exposed sample agrees with a lower refined Na content from neutron powder diffraction data. 109


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Chemistry of Materials The 23Na solid-state NMR spectrum of titanium-substituted P2-phase Na2/3Mn0.8Fe0.1Ti0.1O2 is shown in Figure 4, together

centers with no unpaired electrons (Ti4+), four unpaired electrons (Fe4+, Mn3+), and five unpaired electrons (Fe3+) would be observed. In addition, we should also consider that the presence of two crystallographic sites for sodium in the unit cell of P2 layered structures16 could result, in principle, in two different 23Na solid-state NMR signals. The observation of only one paramagnetic 23Na NMR signal in the spectra is explained by the presence of an averaging effect on the NMR spectrum induced by very fast ion hopping rates in the two-dimensional planes of the layered oxide, which was previously observed in similar Fe + Mn layered compounds.13 The different shifts observed in solid-state NMR spectra of P2-phase Na2/3Mn0.8Fe0.1Ti0.1O2 and P2-phase Na2/3Mn0.8Fe0.2O2 are consequently ascribed to the reduced average unpaired electron density when Ti4+ ions substitute for highly paramagnetic Mn4+ centers in the structure. This result also demonstrates that Ti4+ cations doped into the structure are located at the transition metal layers and not at the sodium layers, as no significant difference in shifts of the two spectra measured in Figure 4 would be expected if Ti were on the Na site. Galvanostatic charge/discharge profile of P2-phase Na2/3Mn0.8Fe0.1Ti0.1O2 from second to 50th cycle is shown in Figure 5a. The pristine electrode delivered a second charge/ discharge capacity of 146.57/144.16 mA·h·g−1 within the voltage range 4.0−2.0 V at C/10, which approximates the utilization of 0.58Na when a theoretical capacity of 245.19 mA· h·g−1 is considered. These specific capacities are comparable to that of similar compounds reported recently,22,23 especially when the significantly faster C-rate and narrow voltage window in the present study are considered. The moisture-exposed electrode delivered a second charge/discharge capacity of 130.40/129.22 mA·h·g−1, which slightly increases to reach the maximum of 134.35 mA·h·g−1 (approximately 10% lower than pristine electrode) as seen in Figure 5b. A slight increase in reversible capacity during the first cycles is probably due to redistribution and stabilization of Na+ and H+ in the interlayer spacing.42−44 It should be noted that the maximum Na reintercalation during discharge for the pristine electrode is close to 0.9Na according to the electrochemical data, which is a theoretical limit due to the existence of 10% Ti4+. This shows remarkable stability of the intercalation process even though all the Mn is reduced to Mn3+, which should induce a large Jahn− Teller distortion at the end of discharge process.45 In addition, the concentration of Na during the charge process up to 4.0 V did not reach beyond 0.35Na, which implies that the P2 → O2 transition does not occur during the deintercalation process.46

Figure 4. 23Na solid-state NMR spectra of Ti4+-doped P2-phase Na2/3Mn0.8Fe0.1Ti0.1O2 and unsubstituted P2-phase Na2/3Mn0.8Fe0.2O2.

with the corresponding spectrum of unsubstituted P2-phase Na2/3Mn0.8Fe0.2O2 for comparison. The unsubstituted compound is used as a reference to study, by NMR, the effect of titanium substitution in the local environment and the dynamics of Na+ ions. The spectra of both compounds are characterized by rather broad and intense signals resonating at very large parts per million (ppm) values and by minor peaks close to 0 ppm. These latter peaks are attributed to the presence of minor diamagnetic impurities in the samples, such as traces of sodium carbonate, which were also previously observed in analogous layered compounds with no influence in the electrochemical performance of the materials.13 The large shifts observed for the main signals in the spectra are induced by the difference in unpaired electron density distribution on the transition metal ions affecting the sodium atom. This effect is known in NMR as “paramagnetic shift” and it is very sensitive to the oxidation states of the transition metal centers and to the structural variations, with an influence on the optimal orbital overlapping required for electron delocalization.39−41 Paramagnetic shifts indicate both changes in the oxidation state of the transition metals during electrochemical charge/discharge, evidenced by ex situ solid-state NMR experiments, and structural changes associated with P2/O3 structures with the same stoichiometry.13 It is worth noting that in a static situation with all Na+ ions “fixed” in the structure, the presence of multiple NMR signals with very different paramagnetic shifts for sodium ions surrounded by different ratios of metallic

Figure 5. Charge/discharge profile of (a) pristine P2-phase Na2/3Mn0.8Fe0.1Ti0.1O2 electrode, (b) moisture-exposed electrode. (c) Plot of capacity retentions and Coulombic efficiency (blue) as a function of cycle number for pristine and moisture-exposed electrode within the voltage range 4.0− 2.0 V at C/10 for 50 cycles. 110


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Figure 6. (a, b) Cyclic voltammograms of (a) pristine and (b) moisture-exposed phases within the voltage window 4.0−2.0 V at a scan rate of 0.1 mV·s−1. (c, d) Rate capability of (c) pristine P2-phase Na2/3Mn0.8Fe0.1Ti0.1O2 electrode and (d) moisture-exposed electrode within the voltage range 4.0−2.0 V. Applied current densities are 24.50 × 10−3 (C/10), 49.00 × 10−3 (C/5), 245.0 × 10−3 (1C), 1.225 (5C), 2.449 (10C), and 12.25 (50C) A·g−1.

Figure 7. Charge/discharge profile of (a) pristine P2-phase Na2/3Mn0.8Fe0.1Ti0.1O2 electrode and (b) moisture-exposed electrode. (c) Plot of capacity retention and Coulombic efficiency as a function of cycle number within the voltage range 4.0−2.0 V at 1C for 300 cycles. The Coulombic efficiency (blue) of moisture-exposed electrode is not shown because of large deviation from 94% to over 100%.

Comparatively, Na reintercalation is limited to ∼0.75Na for the moisture-exposed electrode, presumably due to the presence of protons in the interlayer space. The average voltage and energy density of the pristine electrode are calculated to be 2.77 V and 399.32 W·h·kg−1 versus Na/Na+, respectively, and this energy density approaches that of LiMn2O4, a cathode for LIB.47,48 Within the given voltage range, the charge/discharge curves of both electrodes are smooth compared to the unsubstituted counterpart, without any apparent plateau,21 which is indicative of improved structural stability and a solid solution (secondorder) process of intercalation/deintercalation. Additionally, Fe appears to be electrochemically inactive during the deintercalation/intercalation process because the Mn to Fe redox transition plateau does not appear. As seen in Figure 5c, the pristine electrode exhibited excellent capacity retention by maintaining 95.09% of discharge capacity at the 50th cycle

compared to second cycle (0.1% capacity degradation on each cycle) along with over 98% Coulombic efficiency. The moisture-exposed electrode exhibited even better performance by retaining over 99% of reversible capacity (based on the maximum and 50th discharge capacity) after the initial increase in specific capacity, and the long-term cyclability of the hydrated/protonated phase of similar compounds is reported elsewhere.29−31 In the meantime, the capacity retention plot of Fe-only substituted electrode, P2-phase Na2/3Mn0.8Fe0.2O2, exhibits low Coulombic efficiency less than 80% and capacity retention of 83.78% for 50th cycle (see Figure S4 in Supporting Information). Figure 6a,b shows cyclic voltammograms (CV) of pristine and moisture-exposed electrodes within the voltage range 4.0− 2.0 V at a scan rate of 0.1 mV·s−1. Open-circuit voltage (OCV) of moisture-exposed electrode is slightly higher than that of 111


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Figure 8. In situ XRD patterns of pristine electrode cycled within the voltage range 4.0−2.0 at C/50. Al and Be window peaks are labeled. Regions a and b indicate solid solution mechanisms, and region c includes biphasic mechanism.

during the deintercalation/intercalation process but is rather a complex reaction between electrodes and electrolyte.51 Coulombic efficiency is near 100% for all cycles, while the polarization (18.93−8.18 mV) is almost identical to that of C/ 10. The moisture-exposed electrode, however, degrades faster under the same conditions and maintains only 70% of the discharge capacity of the second cycle after 300 cycles, as seen in Figure 7c. In addition, high polarization of over 800 mV and wide range of Coulombic efficiency from 94% to over 100% during the electrochemical cycling are indicative of a complex degradation mechanism and reduced kinetics due to diffusion hindrance. Due to the low redox potential of Mn3+/4+ reaction, Mn-rich electrode exhibits lower energy density compared to already published high energy density electrodes.52−54 However, the present electrode exhibits excellent capacity retention and rate capability. Figure 8 shows an in situ XRD stacked plot and Bragg diffraction peak evolution for the pristine electrode during the first cycle. Lattice parameters before cycling are calculated to be a = 2.9159(3) Å and c = 11.183(3) Å, which are in good agreement with the values for pristine powder. As evidenced by a lack of formation of new peaks and constant shift in 2θ value of existing peaks, only the solid solution mechanism is observed for the given voltage range during the charge process (region a). All the Bragg diffraction peaks shift without changing in intensity or peak splitting. Such an observation is in good agreement with the smooth cycling profile that exhibits no plateau behavior. Lattice parameter evolution during charge reflects increasing concentration of the smaller Mn4+ by a linearly decreasing a lattice parameter. The initial increase in the c lattice parameter up to 3.53 V is due to stronger adjacent oxide layer repulsion during removal of Na+ from interlayer space (see Figure S6 in Supporting Information). The decrease in the c lattice parameter after 3.53 V, at which the maximum value is reached, is a phenomenon that is observed with similar

pristine electrode, probably because of reduced charge transfer due to Na+/H+ exchange. Other than a peak at 2.51 V (OCV), no peaks relating to redox couple or phase change are observed for the pristine electrode, which is in good agreement with a smooth charge/discharge profile without plateau behavior. The CV curve of the moisture-exposed electrode is somewhat noisy during the first few cycles but becomes smoother in later cycles. This observation is associated with the capacity retention plot, where a slight increase in reversible capacity occurs during the first few cycles. As seen in Figure 6c, the pristine electrode exhibits a remarkable rate capability by maintaining 99.40 mA· h·g−1 discharge capacity at 1C, which is 68.95% of C/10, and such high rate capability indicates fast deintercalation/ intercalation kinetics of the pristine electrode, which also is observed in Ni-substituted P2 and P3 phases.49 The Ragone plot of the pristine electrode (see Figure S5 in Supporting Information) displays typical behavior of a supercapacitor under these experimental conditions, and the energy density and power density are similar to those of a pseudocapacitor.50 Such high rate capability is unique among the Na layered oxides when the bulkiness of the Na+ ion is considered, and further enhancement in rate capability could be possible by reducing the particle size or carbon coating. As seen in Figure 6d, the moisture-exposed electrode exhibits a slightly lower rate capability, especially at 5C or higher rate, which may be attributed to slower kinetics because of interlayer protons acting as Na+ diffusion blocks. Figure 7a,b shows cycling profiles and capacity retention plots of pristine and moisture-exposed electrodes at 1C for 300 cycles. For the pristine electrode, exceptional cyclability is achieved by maintaining 87.17 mA·h·g−1 at the 300th cycle, which is 87.70% capacity retention from the second cycle (0.041% capacity degradation on each cycle). Less degradation per cycle at 1C compared to C/10 indicates that the capacity degradation mechanism is not solely due to material instability 112


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Chemistry of Materials Li ion-containing layered oxides, where the c lattice parameter reaches a maximum value at approximately 70% state of charge and is possibly explained by ordering of the sodium vacancy concentration. At the end of the charge process, even at charge voltage up to 4.3 V, no peaks relating to the P2 → O2 phase transition around at 15.7° 2θ13 are observed (see Figure S7 in Supporting Information), which is in good agreement with the electrochemical charge/discharge and CV profile. A stoichiometric concentration of Na0.37Mn3+0.27Mn4+0.53Fe3+0.1Ti4+0.1O2 at the end of charge (4.0 V) based on the electrochemical data implies a greatly reduced Jahn−Teller distortion compared to the pristine state of Na0.67Mn3+0.57Mn4+0.23Fe3+0.1Ti4+0.1O2. During the subsequent discharge, mostly a solid solution mechanism is observed (region b) where the shift in the Bragg diffraction peaks follows the exact opposite trend to that of the charge process. Lattice parameter evolution follows the opposite of the charge process as well, which indicates a highly reversible structural evolution during the charge/discharge process. However, a biphasic mechanism appears at or below 2.2 V during the discharge process (region c), indicated by peak splitting of the (002) and (004) Bragg diffraction peaks along with peak intensity reduction. The appearance of a small new phase is attributed to increasing distortion caused by the increasing concentration of Mn3+. The new phase is identified as orthorhombic P′2 phase with Cmcm space-group symmetry, and lattice parameters were calculated to be a = 2.896(1) Å, b = 5.376(4) Å, and c = 10.959(5) Å. Upon formation of the new P′2 phase, the evolution of c lattice parameter of the original phase exhibits a constant value while that of P′2 phase decreases, which is consistent with the contracting trend of interlayer space due to increasing concentration of Na. The a lattice parameters of original and new phase exhibit the same flat behavior, while the b lattice parameter of new phase shows an increasing trend, in which the expanding trend of oxide layer continues along the b-direction (see Figure S6 in Supporting Information). Therefore, the lattice parameter evolution, in particular the stacking axis, reflects not only the formation of new phase but also an approximately linear trend as Na concentration increases. At the end of discharge, approximately 0.9Na is reintercalated on the basis of the electrochemical data, which is the theoretical Na intercalation limit where the oxidation state of Mn is entirely 3+. The fact that only a small region of a distorted phase appeared and the discharge process reached a theoretical limit of Na concentration indicates that the effect of Jahn−Teller distortion is reduced by relatively small substitutions of Fe and Ti. Therefore, the structural stability of the current material is vastly improved compared to the unsubstituted Na0.66MnO2, in which the ex situ XRD study concluded that the original structure was not maintained during electrochemical cycling due to complex phase transitions throughout the process.21 Regardless, the biphasic mechanism in the present compound is reversible, as seen in the comparison between XRD patterns of pristine electrode and the cell stopped at the OCV during the second charge process. All the Bragg peaks overlap with each other and the P′2-phase peaks are completely transformed into the original P2-phase peaks (see Figure S7 in Supporting Information). In order to show the inactivity of the Fe3+/4+ redox reaction during cycling, Mössbauer spectroscopy has been applied to the electrode charged up to 4.3 V. As seen in the spectrum (see Figure S8 in Supporting Information), only two contributions are observed, both of them with isomer shifts (IS) in the Fe3+ range of values. Values of the quadrupolar splitting (QS)

confirm that the crystallographic and electronic surroundings of both Fe3+ are significantly different. Fe3+(1) is located in O3NaFeO2-type phase, and the very asymmetric Fe3+(2) belongs to monoclinic NaFeO2-type phase. The weight of both subspectra is the same, indicating that there are equal amounts of Fe3+(1) and Fe3+(2) in the sample (error = 4−5%). Therefore, the Fe3+/4+ redox reaction is not active up to 4.3 V charge process.55,56

4. DISCUSSION Typically, Na containing layered oxide, NaMTO2, crystallizes as P2 or O3 phase based on the concentration of Na in the pristine state, where near or fully sodiated phases typically stabilize as the O3 phase, while lower Na concentration phases (∼0.66Na) stabilize in the P2 phase.10,14 Also, recent publications show that P2 and O3 phases of the same Na concentration could be prepared by changing the precursors and synthesis conditions.13 Mn-rich layered oxides are versatile in terms of synthetic processes due to the existence of multiple oxidation states that are heavily influenced by atmospheric conditions and charge balance of the desired stoichiometry. However, the high-spin-state Mn3+, octahedrally coordinated with O2−,57,58 induces a strong Jahn−Teller distortion during the intercalation process and, therefore, imposes strong crystal strain. In fact, pristine O3-phase NaMnO2 crystallizes as a highly distorted monoclinic phase19 while P2-phase Na0.6MnO2 crystallizes in hexagonal symmetry,21 where the P2 phase signals a decrease in distortion by decreasing the concentration of Mn3+ and increasing the concentration of Mn4+ (t2g3eg0). Thus, the rapid capacity degradation of NaxMnO2 electrodes despite the large initial specific capacity can be attributed to structural strain (stress/relief mechanism) due to the continuous decreasing or increasing concentration of Mn3+ during cycling.21 XRD and powder neutron diffraction data analyses confirm that pristine Na2/3Mn0.8Fe0.1Ti0.1O2 adopts a hexagonal structure in P63/mmc symmetry with no apparent impurity peaks associated with starting precursors or byproducts. Upon exposure to moisture for an extensive amount of time, the hexagonal P63/mmc symmetry still persists, and no impurity peaks are observed. However, the main phase of moistureexposed sample does not appear to be a hydrated phase but a protonated phase with a slightly expanded lattice in the c direction. In addition, the concentrations of Na, Mn, Fe, and Ti, investigated by EDS, ICP, and Rietveld analysis for both electrodes, show similar trends. 23Na NMR spectra display peak shifts associated with substitution of Ti4+ (no paramagnetic contribution) into the transition metal oxide layer, which confirms the successful substitution of Ti4+ into octahedral coordination. The electrochemical performance comparison between C/10 and 1C for both electrodes yields an interesting aspect. At the slow rate of C/10, the deintercalation/intercalation process is sufficiently slow for Na+ to diffuse around any protons, which could provide an anchoring effect for the layered structure, and therefore, the capacity retention for moisture-exposed electrode is better than that of pristine electrode. On the other hand, the deintercalation/intercalation process is so fast at 1C that the protons act more as diffusion obstacles than an anchoring element, and thus the electrochemical performance is greatly reduced. Essentially, from the crystallographic structural data, the variation in electrochemical performance between pristine and moisture-exposed samples can attributed to a combination 113


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NMR confirms the proposed structure of P2-phase Na2/3Mn0.8Fe0.1Ti0.1O2, where Ti4+ ions are integrated into the transition metal oxide layer, and shows that Na+ ions are undergoing very fast dynamics in the two-dimensional sodium layer of the structure. The pristine electrode exhibits stable cycling by maintaining 95.09% of discharge capacity after 50 cycles at C/10 and 87.70% after 300 cycles at 1C. The moisture-exposed electrode exhibits slightly lower reversible capacity, lower rate capability, and more profound capacity degradation at 1C, presumably due to the reduced kinetics resulting from Na+/H+(D+) exchange. However, superior capacity retention of over 99% after 50 cycles of the moisture-exposed phase at C/10 indicates an additional ratedependent stability, presumably due to surface and particle morphology modification during H2O (D2O) treatment. Although the energy density of the present material may be lower than those of already reported ones due to the lower redox potential of Mn3+/4+ reaction, the capacity retention and rate capability are among the best. Structural stability of the pristine electrode is demonstrated by in situ XRD, where only a solid solution mechanism is found during the charge and a small biphasic mechanism is observed below 2.2 V during the discharge. Such improved structural stability at a low level of substitution is attributed to the dilution effect of the Jahn− Teller distortion and strengthening of structural stability of the oxide layer by introducing electrochemically inactive elements like Fe and Ti. In addition, high capacity retention and structural stabilization at such low levels of substitution implies that full utilization of maximum energy density with long-term stability could be possible for a NIB cathode.

of the following structural facets: on average, smaller particle size distribution (as determined by neutron powder diffraction data peak shapes) for D2O sample relative to pristine sample, and likely some Na+/D+ exchange leading to possible formation of hydroxyl groups in the D2O sample. One or a concerted combination of such structural parameters may stabilize the moisture-exposed samples to longer term cycling. The electrochemical process is solely based on the Mn3+/4+ redox process because Ti4+ is electrochemically inactive in this voltage range and the Fe3+/4+ redox plateau does not appear up to 4.0 V. During the discharge process, Mn should be reduced toward the 3+ oxidation state where high-spin Mn3+ induces a strong Jahn−Teller distortion, and the opposite is true during the charge process due to the increase in concentration of highspin Mn4+. Thus, the rapid capacity fade as well as multiple plateau behavior of the unsubstituted compound during electrochemical cycling could be largely attributed to structural strain. On the other hand, the partially substituted Na2/3Mn0.8Fe0.1Ti0.1O2 compound exhibits vastly enhanced electrochemical performance, including smoothness of the potential profile and excellent capacity retention. A previous report speculated that the enhanced electrochemical performance could be due to increasing concentration of the nonJahn−Teller distorted ion, of Mn4+ in the oxide layer.22 However, the average oxidation state of Mn3.29+ in the present case contains higher concentrations of Mn3+ than the unsubstituted compound of Mn3.34+,21 if it is assumed that oxidation states of Fe3+ and Ti4+ are stable in the pristine state. In addition, structural strain during electrochemical cycling should be consistent regardless of the concentration of Mn4+ in pristine state, because intercalation of Na should increase the concentration of Mn3+. It has been well characterized that introducing low-spin-state Mn3+ (t2g4eg0) at high pressure or at low temperature, effectively inducing a high-to-low spin transition of Mn, could reduce Jahn−Teller distortion by a dilution effect.59−61 Furthermore, investigation of the substitution effect of LiMn1−xCrxO2 (0 ≤ x ≤ 0.6) has demonstrated that substitution of Cr, which has a larger electronegativity than Mn, can indeed dilute the cooperative Jahn−Teller distortion.62−65 Consequently, more ordered or less distorted crystals could be synthesized when the concentration of Cr increases. Therefore, the stable electrochemistry attained through a small substitution of Fe and Ti is attributed in part to the dilution effect of Jahn−Teller distortion. In addition, Ti substitution into the oxide layer is known to improve the structural stability of layered oxides.46,66 The electrochemical performance comparison between Fe-only substituted and cosubstituted electrodes further solidifies the importance of Ti in the oxide layer to stabilize the structure. The cooperative effect of the large substitution of various transition metals into Mn-containing layered oxides has been demonstrated either to attain structural stability at the expense of specific capacity or to gain a large specific capacity at the expense of structural stability.50,67,68 Attaining vastly improved structural stability while retaining high specific capacity at low levels of substitution in the present study illustrates the possibility of full utilization of maximum energy density with long-term stability for a layered oxide as a NIB cathode.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.5b03276. Six figures and two tables showing XRD pattern matching of moisture-exposed electrode, Rietveld refined crystallographic parameters from neutron powder diffraction data, elemental analysis, capacity retention and Coulombic efficiency of electrode, Ragone plot, lattice parameter evolution of pristine electrode during in situ XRD, ex situ XRD of pristine and cycled electrode, and Mössbauer spectrum of cycled electrode (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*(M.H.H.) Telephone +34 945297108; e-mail: mhan@ cicenergigune.com. *(T.R.) Fax +34 946013500; telephone +34 946012458; e-mail [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was financially supported by LINABATT project from Ministerio de Economiá Competitividad (ENE201344330-R) and Gobierno Vasco (ETORTEK CIC ENERGIGUNE 10). We express special thanks to Nuria Gomez for ICP elemental analysis, Dr. Jon Ajuria for SEM work, and Egoitz Martin for XRD experiments. Also, Dr. Daniel Carrizo is acknowledged for fruitful discussion regarding supercapacitors.

5. CONCLUSIONS A stable, high-performance cathode material, P2-phase Na2/3Mn0.8Fe0.1Ti0.1O2, has been synthesized and shows remarkable electrochemical performance. 23Na solid-state 114


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