3O2, a Potential

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Cite This: ACS Appl. Mater. Interfaces 2019, 11, 22051−22066

Comprehensive Review of P2-Type Na2/3Ni1/3Mn2/3O2, a Potential Cathode for Practical Application of Na-Ion Batteries Jiaolong Zhang,† Wenhui Wang,*,‡ Wei Wang,† Shuwei Wang,† and Baohua Li*,† †

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Engineering Laboratory for the Next Generation Power and Energy Storage Batteries, Graduate School at Shenzhen, Tsinghua University, Shenzhen 518055, China ‡ Shenzhen Key Laboratory of Organic Pollution Prevention and Control, Environmental Science and Engineering Research Center, Harbin Institute of Technology, Shenzhen 518055, China ABSTRACT: P2-type Na2/3Ni1/3Mn2/3O2 is a promising cathode material for practical applications in Na-ion batteries, due to its high energy density, high volumetric capacity, excellent Na ion conductivity, ease of synthesis, and good stability in air. Yet, it is subject to structural rearrangements on charging to high voltage/ low Na content and Na+/vacancy ordering transitions, which lead to poor reversibility and dramatic capacity decay upon cycling. In this Review, we present the latest advances related to Na2/3Ni1/3Mn2/3O2, with a main focus on strategies to stabilize the structural framework and improve the electrochemical properties. Practical issues and challenges are also proposed on the basis of current research status and progress. KEYWORDS: Na-ion batteries, cathode, P2-Na2/3Ni1/3Mn2/3O2, cation substitution, surface modification, structural transformation, Na+/vacancy ordering

1. INTRODUCTION The energy storage system (ESS) is undoubtedly an essential component in the implementation of intermittent renewable energy in the power grid and the realization of smart grid management by balancing electric energy between peak and off-peak periods.1−4 The most important parameters of an ESS for stationary grid application are low operation cost, long life, high safety, and large volumetric energy density.5,6 As the most successful and sophisticated rechargeable battery developed so far, the Li-ion battery (LIB) is preferred because it satisfies the performance requirements including long life and large volumetric energy density.7−10 However, with the rapid increase in demand for batteries in portable electronics, electric vehicles, etc., concerns about the scarcity of lithium reserves on Earth arise.11−13 In contrast to the limited lithium resource, sodium is earth-abundant, distributed worldwide, and cost-effective. Hence, the Na-ion battery (NIB), which mainly relies on the sodium resource, is considered to be a promising alternative to LIB regarding large-scale battery usage.14−17 Moreover, the Cu current collector in the negative electrode can be replaced with a much cheaper Al current collector in a NIB because sodium does not alloy with Al.5,18,19 Thus, the NIB has attracted much attention since 2012.20−22 An estimation of the relative costs of different components in the NIB shows that the cathode material accounts for the highest proportion, at 32.4%, followed by the separator and anode, with proportions of 14% and 13%, respectively.23 Besides, the cathode is one of the key factors determining the electrochemical performance of the NIB. As a consequence, © 2019 American Chemical Society

exploration of cathode materials with high energy density and good stability is an indispensable part of NIB studies. Among the various cathode choices, polyanionic compounds and layered transition metal oxides are the most widely investigated materials. Polyanionic compounds exhibit high average voltage because of the inductive effect of the polyanion group. However, the presence of heavy polyanion groups reduces the gravimetric capacity of electrode materials.12 More importantly, the low density of polyanionic structures leads to low volumetric energy density.6 Meanwhile, layered transition metal oxides typically deliver high reversible capacity. In addition, their compact structural framework makes them appropriate for grid applications where a high volumetric energy density is desirable.6,24 The structure of layered transition metal oxides depends on oxygen stacking ordering and occupation sites of Na ions. The most prevailing structures are P2 and O3, which were described by C. Delmas in the early 1980s.25 The letters “P” and “O” represent the coordination environment of alkali metal, “prismatic” and “octahedral”, respectively. The numbers “2” and “3” are the number of transition metal (TM) layers in the stacking repeat unit. A schematic illustration of crystal structures of P2 and O3 NaxMeO2 is depicted in Figure 1. P2type NaxMeO2 consists of sheets of MeO6 octahedra with Na ions sandwiched between them. All the Na ions reside in Received: March 4, 2019 Accepted: May 28, 2019 Published: May 28, 2019 22051

DOI: 10.1021/acsami.9b03937 ACS Appl. Mater. Interfaces 2019, 11, 22051−22066

Review

ACS Applied Materials & Interfaces

immersed in water for 1 year.32−34 This allows ease of handling of the material in practical production. (4) Large-scale production of phase-pure P2-Na2/3Ni1/3Mn2/3O2 can be realized via a conditional sintering process followed by natural cooling, which is strikingly different from the other P2-type oxide cathode.35 Most P2-type oxides have to be quenched instantly after the sintering procedure at high temperature (typically 850−950 °C), as the P2 phase is not stable during the slow cooling process;36−40 the quench operation is actually hazardous and difficult to handle. Based on the above properties, we believe Na2/3Ni1/3Mn2/3O2 should be one of the most promising cathode candidates for the practical application of high-energy NIB. In this Review, we present recent developments related to Na2/3Ni1/3Mn2/3O2 with a focus on modification methods to improve the sodium storage performance of Na2/3Ni1/3Mn2/3O2. Critical issues that need to be addressed in future are also discussed in this Review.

Figure 1. Schematic of crystal structures of O3-type and P2-type NaxMeO2. Reprinted with permission from ref 31. Copyright 2015 Electrochemical Society.

prismatic sites, and oxygens are packed stacked in an “ABBA” manner. In the O3 structure, the Na ions are located at octahedral sites between TM layers, and the oxygen stacking arrangement is “ABCABC”. The P2 phase commonly shows a P2→O2 transition upon electrochemical cycling, as the P2-toP3 (or O3) phase transitions require breaking of Me−O bonds and occur only at elevated temperatures. In contrast, O3-type NaxMeO2 usually undergoes complex phase transitions of O3→P3→O′3→P′3 upon electrochemical cycling, which take place at room temperature without breaking Me−O bonds. Therefore, the structure of the P2 phase is more stable than that of the O3 phase during the charge/discharge process.6,11,26 Moreover, the P2 structure provides more spacious diffusion paths for Na ions.26,27 As a result, the P2 phase exhibits much better cycle stability and rate performance.26,28−30 Na2/3Ni1/3Mn2/3O2 attracts great attention among the P2type materials in view of its electrochemical performance and material handling. (1) Na2/3Ni1/3Mn2/3O2 exhibits a high operating voltage (i.e., ∼3.8 V) and high capacity (i.e., ∼173 mAh/g), which in return enable it to have the highest energy density among layered transition metal oxide cathodes when the full cell is assembled.1,6,32 (2) A majority of P2-type oxides are sodium deficient, while the amount of sodium ions in Na2/3Ni1/3Mn2/3O2 is sufficient for electrochemical cycling, which is a significant benefit for achieving a high-energy full cell, as no extra components (e.g., sodium-rich sacrificial source) are needed to compensate for a sodium deficiency during the first charging/discharging process. (3) Unlike the hygroscopic property of most sodium-based layered oxides, Na2/3Ni1/3Mn2/3O2 is stable in ambient atmosphere, and the crystal structure can be well preserved even after being

2. OVERVIEW OF P2-NA2/3NI1/3MN2/3O2 The sodium storage behavior of P2-Na2/3Ni1/3Mn2/3O2 was first investigated by J. R. Dahn’s group in 2001.41 A Na2/3Ni1/3Mn2/3O2 electrode is able to deliver a reversible capacity of ∼160 mAh/g over 2.0−4.5 V at a current density of 5 mA/g, which is close to the theoretical capacity of ∼173 mAh/g. Hence, it is speculated that almost all the Na ions can be reversibly extracted from/inserted into Na2/3Ni1/3Mn2/3O2. During the charging process, NaxNi1/3Mn2/3O2 undergoes a phase transformation, accompanied by the exaction of Na+. Pristine P2 structure persists until x decreases to ∼1/3 (charges to ∼4.2 V), at which point O2-type stacking faults start to appear due to gliding of the oxygen layers. A further decrease of Na content to x < 1/3 leads to coexistence of the P2 and O2 phases. The P2−O2 transformation during Na+ extraction is consistent with expectation based on the calculated formation energy. 42 During the subsequent discharging process, the re-intercalated Na+ could gradually eliminate the stacking faults and enable re-formation of the pristine P2 structure. Evolution of lattice parameters upon charging/discharging was also studied. Upon charging, the c axis expands when x > 1/3 due to the increased electrostatic repulsive force between adjacent TM layers at lower Na content; however, c begins to decrease because of the closepacked arrangements of O atoms as well as the newly emerged O2-type phase with smaller c when x < 1/3. The a axis, which is dominated by the M−M distance, keeps on shrinking upon charging as a result of oxidation of the Ni ions. A reverse process occurs upon discharging Na2/3Ni1/3Mn2/3O2.

Figure 2. Charge/discharge curves of a Na2/3Ni1/3Mn2/3O2/Na cell over different voltage ranges at a current density of 0.1 C: (a) 2.0−4.5, (b) 2.0−4.0, and (c) 1.6−3.8 V. Reprinted with permission from ref 43. Copyright 2013 Elsevier. 22052


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is to build an artificial physical barrier. Hence, surface modification has been carried out for Na2/3Ni1/3Mn2/3O2 material. Moreover, selection of an appropriate binder was also demonstrated as an alternative way to improve the sodium storage performance of Na2/3Ni1/3Mn2/3O2, as it helps to stabilize the electrode surface and suppress electrolyte decomposition. 3.1. Inert Cation Substitution. Li and Mg substitution for Na2/3Ni1/3Mn2/3O2 has been comprehensively investigated. It is generally acknowledged that the Li dopants preferentially occupy TM layers, as the sites in Na layers are too large for Li to exist stably. The presence of Li leads to a phase-transitionfree characteristic of the Na2/3Ni1/3Mn2/3O2 cathode, thus enhancing the structural stability.49−51 As can be seen from Figure 3a, the stepwise voltage profiles transform to smooth

Figure 2a shows the typical charge/discharge curves over 2.0−4.5 V, and the long plateau at ∼4.2 V is identified as a biphasic region.43 The capacity rooted from the 4.2 V plateau decays quickly upon cycling, accompanied by a sharp increase in polarization. As a consequence, the discharge capacity drops from ∼140 mAh/g to ∼40 mAh/g from the first cycle to the 100th cycle. These results suggest that the P2−O2 structural transformation could result in severe structural degradation and capacity loss of Na2/3Ni1/3Mn2/3O2, although the structural transformation is initially reversible. Na2/3Ni1/3Mn2/3O2 also suffers phase transitions between different Na+/vacancy ordered patterns formed at particular Na stoichiometries, as reflected by the multiple voltage steps in charge/discharge curves. Na+/vacancy ordering is prevalent in Na-based layered oxides due to strong Na+−Na+ repulsion in the Na layer and charge ordering in the TM layer.6 The formation of ordered superstructures is detrimental to electrochemical properties because the superstructures lower the Na-ion diffusion coefficient and reduce the dimensionality of Na ion transport.44,45 Electrolyte decomposition at high voltage is another reason for the poor cycle performance of Na2/3Ni1/3Mn2/3O2. Confining the upper cutoff voltage has been proposed as one strategy to improve the long-term cycle stability of Na2/3Ni1/3Mn2/3O2, as the P2−O2 structural transformation at ∼4.2 V and electrolyte decomposition at high voltage are the major causes of its capacity fading.42,43 As shown in Figure 2b, a stable capacity of ∼84 mAh/g is retained at the 100th cycle in a limited voltage range of 2.0−4.0 V, which is about twice that for cycling over 2.0−4.5 V. On the other hand, voltage confinement inevitably causes a sacrifice in capacity. Therefore, reducing the lower cutoff voltage was adopted in an attempt to gain a higher capacity. When cycling was done over 1.6−3.8 V, the first discharge capacity reached ∼135 mAh/g, which is comparable with the capacity obtained when cycling from 2.0 to 4.5 V (see Figure 2c). Only Ni2+ is oxidized during charging, while both Nin+ and Mn4+ are reduced during discharging to 1.6 V, which contributes to the high discharge capacity.46 However, Mn4+/Mn3+ redox also brings about unsatisfactory cycle stability due to the Jahn−Teller effect of Mn3+ and Mn dissolution into the electrolyte. In addition, the energy density during cycling over 1.6−3.8 V is much lower than that during cycling over 2.0−4.5 V, despite their similar capacity, due to the much lower operating voltage when cycling over 1.6−3.8 V. Based on the above discussion, the electrochemical performance of Na2/3 Ni1/3Mn 2/3 O 2 can be enhanced by (1) suppressing P2−O2 structural transformation, (2) eliminating Na+/vacancy ordering, (3) avoiding/reducing participation of the Mn4+/Mn3+ redox couple, and (4) inhibiting electrolyte decomposition.

Figure 3. (a) Selected voltage profiles of Na0.80[Li0.12Ni0.22Mn0.66]O2 over 2.0−4.4 V at 0.1 C (1 C = 118 mAh/g); inset table shows the distribution of Li ions in TM and Na layers at different stages during the first cycle and fifth cycle based on NMR spectra. (b) In situ SXRD of Na0.80[Li0.12Ni0.22Mn0.66]O2 during the first charge to 4.4 V (peaks marked by “*” are reflections from Al current collector). Reprinted with permission from ref 50. Copyright 2014 American Chemical Society.

charge/discharge curves after the introduction of Li, indicating that the reaction mechanism changes from a series of bi-phase transition reactions to solid-solution-like behavior. In addition, Li in TM layers stabilizes the P2 structure over a wide range of Na concentrations. P2 structure of Na0.80[Li0.12Ni0.22Mn0.66]O2 can be maintained throughout the charge/discharge process over 2.0−4.4 V with only variation in the lattice parameters (see Figure 3b), which is strikingly different from the P2−O2 transformation occurring in Li-free sample above 4.0 V. The function of Li in stabilizing the structure is expected to be closely correlated with the migration of Li from the TM layer to the Na layer at a high charging voltage (see the inset table in Figure 3a). The electropositive Li ions that migrate to Na

3. APPROACHES TO IMPROVE SODIUM STORAGE PERFORMANCE OF P2-NA2/3NI1/3MN2/3O2 As mentioned earlier, Na+/vacancy ordering is an important reason for the capacity fading of Na2/3Ni1/3Mn2/3O2. It is suggested that introduction of other transition metal ions with similar ionic radii and different Fermi levels could prevent Na+/vacancy ordering by breaking charge ordering in the TM layer.6,47,48 Therefore, a variety of dopants, including inert and active cations, have been explored for Na2/3Ni1/3Mn2/3O2. Electrolyte decomposition at high voltage also deteriorates the cycle stability of Na2/3Ni1/3Mn2/3O2. An effective way to impede the side reactions between the cathode and electrolyte 22053


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Figure 4. Charge/discharge curves of (a) Na0.67Mg0.05Ni0.25Mn0.7O2 and (b) Na0.67Mg0.1Ni0.2Mn0.7O2 over the voltage range of 2.0−4.5 V at 12 mA/g. Charge/discharge curve at 2.0−4.5 V and corresponding in situ XRD of Na0.67Mg0.1Ni0.2Mn0.7O2 are shown in (c) and (d), respectively. Reprinted with permission from ref 54. Copyright 2016 American Chemical Society.

basis of a redox couple of Ni2+/Ni4+, and the extra capacity is suspected to originate from oxygen redox reaction. This phenomenon becomes pronounced when Mg content reaches 20 mol%. For example, extra charge capacity of ∼38 mAh/g is obtained for Na2/3Ni1/3‑xMgxMn2/3O2 (x = 0.2).57 Ti-doped Na2/3Ni1/3Mn2/3O2 (i.e. Na2/3Ni1/3Mn2/3‑xTixO2(0 ≤ x ≤ 2/3)) materials were first studied by Komaba’s group.58 It is found that Ti doping helps to improve the cycle stability of Na2/3Ni1/3Mn2/3O2 via suppressing Na+/vacancy ordering and reducing volume changes upon Na de-intercalation/intercalation, although doping with electrochemically inactive Ti results in a decrease in the initial reversible capacity. As shown in Figure 5a,b, the optimal Ti-doped material, i.e., Na2/3Ni1/3Mn1/2Ti1/6O2, delivers an initial discharge capacity of 127 mAh/g with capacity retention of 88% for 20 cycles over 2.5−4.5 V at 12.1 mA/g. Volume contraction of the crystal unit cell for Ti-doped samples at fully charged states is 12−13%, which is much smaller than the value of 23.1% for the Ti-free sample. It is also shown that an improved capacity of ∼120 mAh/g can be retained from Na2/3Ni1/3Mn1/2Ti1/6O2 after 20 cycles under constant current−constant voltage (CCCV) mode over 2.5−4.35 V, compared to that under constant current (CC) mode over 2.5−4.5 V (Figure 5c). The improvement is most likely due to the reduction in the upper cutoff voltage from 4.5 to 4.35 V. After cycling over 2.5− 4.35 V under CC-CV mode, the voltage plateau at ∼4.2 V is well maintained, suggesting that the structural transformation at high voltage is highly reversible. P. F. Wang et al.59 carried out a more systematic study on Na2/3Ni1/3Mn1/3Ti1/3O2. As shown in Figure 5d, Na2/3Ni1/3Mn1/3Ti1/3O2 delivers a reversible capacity of ∼88 mAh/g with smooth voltage profiles when cycled between 2.5 and 4.15 V at 0.1 C. No obvious capacity decay and voltage drop are observed after 100 cycles. At 1 C, Na2/3Ni1/3Mn1/3Ti1/3O2 displays an excellent capacity retention of ∼83.9%, even after 500 cycles (Figure 5f). Moreover, Na2/3Ni1/3Mn1/3Ti1/3O2 could still deliver ∼87.9%

layers at high voltage are able to hold the adjacent TM layers together, thus inhibiting oxygen layer gliding and phase transformation.6 As a result, Li doping largely suppresses the capacity loss upon cyclingmore than 90% capacity can be retained at the end of 50th cycle. Mg ions could also function as structural stabilizers for Na2/3Ni1/3Mn2/3O2, thereby improving the cycle life. P. F. Wang et al.52 first reported that 5 mol% Mg doping in the TM layer is sufficient to suppress P2−O2 structural transformation of Na2/3Ni1/3Mn2/3O2. After Na0.67Mn0.67Ni0.28Mg0.05O2 was charged to 4.35 V, all diffraction peaks could be indexed as P2 phase with many Na vacancies.53 In addition, the P2 structure of Na0.67Mn0.67Ni0.28Mg0.05O2 was retained even after 100 cycles over 2.5−4.35 V despite of its reduced crystallinity. Later, G. Singh et al.54 clearly revealed that Na0.67Ni0.2Mg0.1Mn0.7O2 undergoes reversible P2−OP4 transformation in a voltage range of 2.0−4.5 V (see Figure 4d). The OP4 phase has an intergrowth of P2 and O2 structure, which minimizes the drastic P2−O2 transitions via adopting alternate stacking of octahedral and trigonal prismatic Na layers along the c axis direction. Hence, the reversibility of the P2−OP4 transition is superior to that of the P2−O2 transition. A recent publication confirmed those finding using both ab initio simulations and experiments.55 Another interesting finding of the report is that the ordered structures are almost eliminated and a more solid-solution-like reaction proceeds in a Mgsubstituted sample when cycled below 4.0 V. This is because the random distribution of Mg ions effectively suppresses the development of long-range charge ordering within the TM layers, hence inhibiting long-range Na ordering. Mg substitution furthermore increases the Na ion diffusivity as the interlayer space is expanded after Ni ions are replaced with larger Mg ions.56 All these changes lead to a prominent enhancement in capacity retention (see Figure 4a,b). It is also worth mentioning that Mg-substituted samples deliver extra capacity compared to the theoretical capacity calculated on the 22054


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ascribed to the following reasons: (1) Zn doping promotes the reversible P2−O2−P2 transformation upon charging/ discharging process via suppressing the formation of a shortrange disordered “Z” phase at the charged state and an orthorhombic P′2 phase at the discharged state (see Figure 6b,d).53,61,62 (2) Zn doping is able to relieve structural distortion in MnO6 and NiO6 octahedra during cycling, as evidenced by XANES. (3) Zn doping can effectively inhibit the irreversible migration of Ni2+ from octahedral sites in the TM layer to tetrahedral sites in the Na layer at the fully charged state, which would be accompanied by a distortion of the oxygen lattice for pristine Na2/3Ni1/3Mn1/2O2, thus retarding the structural distortion (see Figure 6e).61 Introduction of Al into the Ni sites of Na2/3Ni1/3Mn2/3O2 can significantly improve its long-term cycle stability.63 Ex situ XRD analyses of Na0.6Ni0.22Al0.11Mn0.66O2 in Figure 7d illustrate that reversible P2−O2−P2 phase transformation can be achieved when the cell is charged/discharged between 4.6 and 2.0 V. Nevertheless, an orthorhombic P′2 phase is generated if the cell is further discharged to 1.5 V, because part of the Mn4+ is reduced to Jahn−Teller effective Mn3+ below 2.0 V, and the formation of P′2 is generally believed to be detrimental to the cycle stability. The authors have also tried to optimize the upper cutoff voltage to achieve both high capacity and high stability. The result shows that 2.0−4.3 V is the optimal voltage range for Na0.6Ni0.22Al0.11Mn0.66O2, showing not only slightly better cycle stability but also a much higher capacity than those obtained over 2.0−4.0 V (see Figure 7b). Under this condition, Na0.6Ni0.22Al0.11Mn0.66O2 delivers an initial discharge capacity of ∼130 mAh/g with an average working voltage of 3.6 V at 20 mA/g. At the end of the 200th cycle, a capacity of ∼102 mAh/g can still be obtained, with retention of ∼79% (see Figure 7c). It is worth mentioning that the doped Al element is not homogeneously distributed and an Al-rich compound may form in Na0.6Ni0.22Al0.11Mn0.66O2. J. X. Shen’s group showed that substitution of Mo for Mn in Na2/3Ni1/3Mn1/2O2 expands the interlayer space and generates a Na2MoO4 impurity.49 There is no noticeable enhancement in cycle stability with Mo doping, while the rate performance of Na 0 . 6 6 Ni 0 . 3 3 Mn 0 . 6 1 Mo 0 . 0 5 O 2 is superior to that of Na2/3Ni1/3Mn1/2O2 due to the increased interlayer distance. The effect of inert cation doping on the capacity and cycle stability of Na0.66Ni0.33Mn0.67O2 is summarized in Figure 8. In general, most of the inactive cations help to enhance cycle performance of Na0.66Ni0.33Mn0.67O2 when charged above 4.2 V, especially for the Al dopant. On one hand, substitution with inert cations usually leads to lower capacity, which suggests that more Na ions are allowed to stay in alkali layers for charge balance. Hence, more Na ions hold the adjacent transition metal layers together up to the end of charging, and structural transformation is thus suppressed. On the other hand, Na+/ vacancy ordering is subdued in substituted materials due to the suppressed charge ordering in the TM layer. 3.2. Active Cation Substitution. One major drawback of inert cation substitution in Ni sites of Na2/3Ni1/3Mn1/2O2 is the loss of capacity due to the reduced amount of Ni2+/Ni4+ redox reaction. In this regard, substitution with electrochemically active cations should be a promising alternative approach to stabilize Na2/3Ni1/3Mn1/2O2, as it can improve the cycle stability of the materials via synergistic effect between different transition metals with minimal sacrifice of reversible capacity. Similar to the effect of inert cation doping, replacing some of the Ni with Cu is also capable of inhibiting Na+/vacancy

Figure 5. (a) Charge/discharge curves and (b) cycle performance of Na2/3Ni1/3Mn2/3‑xTixO2(0 ≤ x ≤ 2/3) at 12.1 mA/g under CC mode. The voltage ranges for Na2/3Ni1/3Mn2/3‑xTixO2 (0 ≤ x ≤ 1/3) and Na2/3Ni1/3Mn2/3‑xTixO2 (x = 2/3) are 2.5−4.5 and 2.5−4.2 V, respectively. (c) Charge/discharge curves of Na2/3Ni1/3Mn1/2Ti1/6O2 under CC-CV mode over 2.5−4.35 V. Reprinted with permission from ref 58. Copyright 2014 Royal Society of Chemistry. (d) Charge/ discharge curves of Na2/3Ni1/3Mn1/3Ti1/3O2 (P2-NaNMT) at different cycles at 0.1 C. (e) Rate performance of Na2/3Ni1/3Mn2/3O2 (P2NaNM) and P2-NaNMT. (f) Cycle performance of P2-NaNM and P2-NaNMT at 1 C. The voltage range is between 2.5 and 4.15 V. Reprinted with permission from ref 59. Copyright 2018 American Association for the Advancement of Science.

and ∼77.5% of its initial capacity at 10 and 20 C, respectively, demonstrating an outstanding rate capability. The effect of Ti substitution on Na2/3Ni1/3Mn2/3O2 can be summarized in two aspects: (1) suppression of Na+ /vacancy ordering by restraining electron delocalization and (2) decreasing the activation energy barrier for Na ion hops between adjacent prismatic sites due to the presence of Na+/vacancy disordering, which is thus beneficial for fast ion diffusion and excellent kinetic properties. Systematic investigation on the effect of Zn doping on Ni site in Na2/3Ni1/3Mn1/2O2 has been conducted by Y. Yang’s group using in situ high-energy X-ray diffraction (HEXRD), ex situ X-ray absorption near edge structure (XANES), and solidstate nuclear magnetic resonance (SS-NMR).60,61 Similar to the effect of Ti substitution, Na0.66Ni0.26Zn0.07Mn0.67O2 shows superior cycle stability at the expense of the initial capacity (see Figure 6a,c). The improvement in capacity retention is 22055


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Figure 6. Charge/discharge curves over 2.2−4.3 V at 12 mA/g and in situ HEXRD patterns during first cycle of Na0.66Ni0.33Mn0.67O2 (a, b) and Na0.66Ni0.26Zn0.07Mn0.67O2 (c, d). (e) Schematic diagram of the structural transformation of Na0.66Ni0.33Mn0.67O2 observed experimentally. Reprinted with permission from ref 61. Copyright 2016 American Chemical Society.

Cu2+ (0.73 Å) is larger than that of Ni2+ (0.69 Å)) and lowering the Na diffusion barrier (rooted from the Jahn−Teller active Cu2+ (t2g6-eg3)).67,68 In addition, the introduction of Cu2+ does not alter the high stability of Na0.67Ni0.3Mn0.7O2 toward air or moisture exposure. For example, the crystal structure of Na2/3CuxNi1/3‑xMn2/3O2 (0 ≤ x ≤ 1/4) shows no obvious change after exposure to air or soaking in water.64,65 Five days’ exposure to air has no adverse impact on the initial charge/discharge performance of Na2/3CuxNi1/3‑xMn2/3O2 (0 ≤ x ≤ 1/4).64 Fe-substituted Na2/3Ni1/3Mn1/2O2 is considered to be a promising candidate for use as a NIB cathode due to the utilization of environmentally benign metals. Phase-pure Na0.67Ni0.15Fe0.20Mn0.65O2 with P2 structure can be synthesized by a sol−gel method. 69 As shown in Figure 10c, Na0.67Ni0.15Fe0.20Mn0.65O2 delivers a high initial discharge capacity of ∼208 mAh/g over 1.5−4.3 V at 13 mA/g because of the joint participation of Ni2+/Ni4+, Fe3+/Fe4+, and Mn3+/ Mn4+ redox couples. Fe doping suppresses Na+/vacancy ordering and renders reversible P2−OP4 phase transition via eliminating the formation of the O2 phase during the charge/ discharge process (Figure 10d), thus resulting in good cycle stability of Na0.67Ni0.15Fe0.20Mn0.65O2 (retains ∼71% of its initial discharge capacity at the 50th cycle). J. Hassoun’s group

ordering, as indicated by the reduced number of peaks in the cyclic voltammogram and smoother charge/discharge curves in Figure 9a,b.64−66 The redox activity of Cu2+/Cu3+ couples in Cu-substituted materials is evidenced by the extra reversible capacity beyond the theoretical capacity estimated by the Ni2+/ Ni4+ redox couple and ex situ XPS analysis. The structural evolution of Na0.67Ni0.3−xCuxMn0.7O2 during the charge/ discharge process is significantly affected by the dopant amount. When x in Na0.67Ni0.3‑xCuxMn0.7O2 is lower than 0.2, the crystalline structure of Cu-doped materials at the fully charged state can be stabilized as intermediate OP4 phase instead of O2 phase, which reduces damage resulting from gliding of layers. 40 More interestingly, when x in Na0.67Ni0.3‑xCuxMn0.7O2 reaches 0.2, the Cu-doped materials can maintain the P2 structure during the whole charge/ discharge process over 2.0−4.5 V (Figure 9d).64,65 The stabilized structure induced by Cu2+ doping in return largely improves the cycle stability of Na 0.67 Ni 0.1 Cu 0.2 Mn 0.7 O 2 compared to that of pristine material. As can be seen from Figure 9c, Na0.67Ni0.1Cu0.2Mn0.7O2 delivers an initial capacity of ∼115 mAh/g with capacity retention of ∼90% after 50 cycles in the working voltage range of 2.0−4.5 V at 17 mA/g. Cu substitution also improves the rate performance via expanding the interlayer distance (because the ionic radius of 22056


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Figure 7. Characterization of Na0.6Ni0.22Al0.11Mn0.66O2. (a) Charge/discharge curves and (b) cycle performance within various voltage ranges at 20 mA/g. (c) Long-term cycling behavior over 2.0−4.3 V at 20 mA/g. (d) Ex situ XRD and SEM analyses of the electrodes at different states of charging during the first cycle performed over 2.5−4.6 V at 20 mA/g. Reprinted with permission from ref 63. Copyright 2017 Royal Society of Chemistry.

10c,d).70 It is worth mentioning that sodium leaching takes place when P2-Na0.67Ni0.15Fe0.20Mn0.65O2 is rinsed in water, as evidenced by the lower Na content and increased lattice parameter c in the rinsed sample compared with the pristine material. Co substitution in Na2/3Ni1/3Mn1/2O2 has also attracted much attention.71−76 Z. Y. Li et al.77 showed that the improved electrochemical properties of Co-doped Na2/3Ni1/3Mn1/2O2 can be explained by the following reasons: (1) Co-doped Na2/3Ni1/3Mn1/2O2 has a more stable structure, which may derive from shrinkage of the TM−O and O−O bond lengths and contraction of the MO6 octahedron with Co substitution. (2) Introduction of Co enlarges the interlayer space, which could facilitate Na ion diffusion. (3) Aliovalent substitution of Co for Ni enhances the electronic conductivity of the material. S. Passerini’s group71 found that water rinsing causes Na ions to leach out from the bulk Na0.63Ni0.22Co0.11Mn0.66O2, thus forming a Na-deficient Na0.45Ni0.22Co0.11Mn0.66O2 phase, due to proton/sodium exchange during the water rinsing process. The rinsed Na0.45Ni0.22Co0.11Mn0.66O2 with lower Na content shows better cycling stability as compared to as-prepared

Figure 8. Comparison of capacity and cycle stability of NaxIyNizMn1‑y‑zO2 (0.60 ≤ x ≤ 0.80, 0 ≤ y ≤ 0.17, 0.20 ≤ z ≤ 0.33) substituted with different inert cations.

also reported that P2-Na0.5[Ni0.23Fe0.13Mn0.63]O2, prepared by a co-precipitation method, shows electrochemical behavior similar to that of P2-Na0.67Ni0.15Fe0.20Mn0.65O2 (see Figure 22057


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Figure 9. (a) CV curves of Na0.67Ni0.3Mn0.7O2 and Na0.67Ni0.1Cu0.2Mn0.7O2 at 1 mV/s. (b, c) Initial charge/discharge curves and cycle performance of Na0.67Ni0.3‑xCuxMn0.7O2 (x = 0, 0.1, 0.2, and 0.3) over 2.0−4.5 V at 0.1 C. (d) In situ XRD patterns collected during charging/discharging of Na0.67Ni0.1Cu0.2Mn0.7O2 over 2.0−4.5 V at 0.1 C. (e) XRD patterns of Na0.67Ni0.1Cu0.2Mn0.7O2 electrodes before and after continuous charge/ discharge processes. Reprinted with permission from ref 65. Copyright 2017 Royal Society of Chemistry.

and 4.4 V at 20 mA/g.78 D. Buchholz et al.79 investigated the water sensitivity of P2/P3-NaxNi0.22Co0.11Mn0.66O2 at different charge states. Their study demonstrated that P2/P3NaxNi0.22Co0.11Mn0.66O2 is very prone to take up water to form a hydrated phase at a low Na content of x < 0.33 (or voltage above 3.6 V), due to the increased distance of neighboring oxygen layers and available vacancies in the Na layer. The effects of substitution with different active cations on capacity and cycle stability are summarized in Figure 12. Cycle stability of Na2/3Ni1/3Mn2/3O2 is improved mainly due to the synergistic effect between different transition metals together with suppressed Na+/vacancy ordering. Among the modified materials, introduction of Co into Na2/3Ni1/3Mn2/3O2 helps to improve the cycle stability dramatically without sacrificing the initial capacity. In addition, rinsing the as-prepared material with water was found to be beneficial for Na storage performance. Nevertheless, rinsing also causes loss of Na, thus resulting in the problem of Na deficiency, which is one of the most critical problems limiting the practical application of most P2-type materials.

Na0.63Ni0.22Co0.11Mn0.66O2. This is because the rinsing process promotes the reversibility of the 4.2 V plateau, which is probably due to the stabilized structure induced by the proton/ sodium exchange. In the voltage range of 2.1−4.3 V, rinsed Na0.45Ni0.22Co0.11Mn0.66O2 is capable of delivering a discharge capacity of ∼119 mAh/g at the 100th cycle at 12 mA/g, corresponding to an excellent capacity retention of ∼85%. Even after interleaved testing at various current densities ranging from 0.2 to 5 C, an extremely high discharge capacity of ∼100 mAh/g can be attained after 275 cycles. On the other hand, the capacity degradation of rinsed Na0.45Ni0.22Co0.11Mn0.66O2 persists, accompanied by the gradual shortening of the 4.2 V plateau upon repeated cycling (see Figure 11c). This phenomenon is associated with the fact that structural transformation to the O2 phase at high voltage continues to occur in Co-doped Na2/3Ni1/3Mn1/2O2. Al substitution has also been employed to further stabilize the Na0.67[Mn0.65Ni0.15Co0.2]O2 layered structure. The resulting Na0.67[Mn0.65Ni0.15Co0.15Al0.05]O2 shows negligible degradation in both capacity and voltage, from the highest value of ∼129 mAh/g to ∼123 mAh/g at the 50th cycle between 2.0 22058


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Figure 10. (a) Cycle performance of Na0.67NixFe0.35‑xMn0.65O2 (x = 0, 0.15) over 1.5−4.3 V at 13 mA/g. (b) Ex situ XRD patterns of Na0.67Ni0.15Fe0.20Mn0.65O2 (x = 0.15) at various charge/discharge depths. Reprinted with permission from ref 69. Copyright 2014 Elsevier. (c) Charge/discharge profiles and (d) cycle performance and Coulombic efficiency of rinsed Na0.5[Ni0.23Fe0.13Mn0.63]O2 over 1.5−4.6 V at 15 mA/g. Reprinted with permission from ref 70. Copyright 2014 John Wiley and Sons.

Figure 11. (a) First and second charge/discharge curves of as-prepared Na0.63Ni0.22Co0.11Mn0.66O2 and rinsed Na0.45Ni0.22Co0.11Mn0.66O2 at 0.1 C (1 C = 120 mA/g). (b) Cycle performance of rinsed Na0.45Ni0.22Co0.11Mn0.66O2 at different current densities. (c) Selected charge/discharge curves of rinsed Na0.45Ni0.22Co0.11Mn0.66O2 at 0.1 C. (d) Charge/discharge curves of rinsed Na0.45Ni0.22Co0.11Mn0.66O2 at different current densities. Voltage cutoff limits: 2.1−4.3 V. Reprinted with permission from ref 71. Copyright 2013 American Chemical Society.

J. Alvarado et al.81 investigated the effect of Al2O3 coating on electrochemical performance and the cathode−electrolyte interface (CEI) of Na2/3Ni1/3Mn2/3O2 by atomic layer deposition (ALD). The thickness of the coating layer was approximately 1 nm based on a TEM image. As can be seen from Figure 13a, typical voltage plateaus of Na2/3Ni1/3Mn2/3O2 persist in Al2O3-coated Na2/3Ni1/3Mn2/3O2 over 2.3−4.5 V,

3.3. Surface Modification. Construction of surface coating layers is another common but effective strategy to stabilize the high-voltage cathode materials, as they minimize the interaction between electrode and electrolyte.80 To date, only a few coating layers have been tested for Na2/3Ni1/3Mn2/3O2. 22059


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barrier becomes a control factor and leads to a decrease of reversible capacity (Figure 13c). Y. Liu et al.82 proposed a wet chemistry method (WCM) to prepare Al2O3-coated Na2/3Ni1/3Mn2/3O2 with a targeted coating amount of 5 wt% of the total mass. Nevertheless, the electrochemical behavior of WCM-Al2O3-Na2/3Ni1/3Mn2/3O2 is quite different from that of the above-mentioned ALDAl 2 O 3 -Na 2/3 Ni 1/3 Mn 2/3 O 2 . First of all, WCM-Al 2 O 3 Na2/3Ni1/3Mn2/3O2 displays relatively smooth charge/discharge curves in Figure 14b, as the voltage plateaus below 3.7 V in Na2/3Ni1/3Mn2/3O2 are replaced by a sloping curve after Al2O3 modification, suggesting that WCM-Al2O3Na2/3Ni1/3Mn2/3O2 undergoes a different Na storage mechanism and is more inclined to a solid-solution-like behavior. In general, pure coating with an inactive layer cannot change the reaction mechanism. Therefore, we infer that at least some of the Al is doped into Na2/3Ni1/3Mn2/3O2. It is possible for doping to occur during the high-temperature annealing (650 °C, 10 h) process during coating of WCM-Al 2 O 3 Na2/3Ni1/3Mn2/3O2. The hypothesis of Al doping in WCMAl2O3-Na2/3Ni1/3Mn2/3O2 is also supported by the following two points: First, WCM-Al2O3-Na2/3Ni1/3Mn2/3O2 exhibits significantly improved cycle stability compared with Na2/3Ni1/3Mn2/3O2 (Figure 14c), rather than the small improvement made by the ALD-coated sample in J. Alvarado’s work,81 where the doping effect was eliminated. Second, WCM-Al2O3-Na2/3Ni1/3Mn2/3O2 delivers not only similar smooth charge/discharge curves but also cycle stability comparable to that of the Al-doped sample reported by I. Hasa et al.63 Except for suppressing Na+/vacancy ordering, the WCM modification also protects the crystal structure by inhibiting particle exfoliation upon repeated cycling, as can be seen from Figure 14e,f.

Figure 12. Comparison of capacity and cycle stability of NaxAyNizMn1‑y‑zO2 (0.45 ≤ x ≤ 0.67, 0 ≤ y ≤ 0.20, 0.10 ≤ z ≤ 0.33) substituted with different active cations.

indicating that the coating layer does not alter the Na ion intercalation/de-intercalation behavior. The insulating nature of the Al2O3 layer has a negligible effect on capacity at a current density of 0.05 C, as sufficient time is allowed for Na ions to migrate through the coating layer at low current density. On the other hand, Al2O3 coating can reduce the electrolyte decomposition, thus avoiding the formation of overly thick SEI and facilitating high reversible capacity. As shown in Figure 13b, Al2O3-coated sample retains a capacity of ∼77.4 mAh/g at the end of the 100th cycle, which is higher than ∼52.0 mAh/g obtained from the uncoated sample. The improved electrochemical performance should originate from the reduced particle exfoliation and the formation of better CEI enabled by the Al2O3 coating (Figure 13d). At a high current density of 1 C, the time for Na ions’ migration is limited. Therefore, the kinetic hindrance induced by the Al2O3

Figure 13. (a) First charge/discharge curves, (b) cycle performance at 0.05 C, (c) rate performance, and (d) Nyquist plots and SEM images at the end of the 100th cycle for Na2/3Ni1/3Mn2/3O2 and ALD-Al2O3-Na2/3Ni1/3Mn2/3O2 samples. The voltage range is 2.3−4.5 V. Reprinted with permission from ref 81. Copyright 2017 American Chemical Society. 22060


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Figure 14. Electrochemical, morphological, and ex situ XRD study of Na2/3Ni1/3Mn2/3O2 and WCM-Al2O3-Na2/3Ni1/3Mn2/3O2: (a) charge/ discharge curves of Na2/3Ni1/3Mn2/3O2, (b) charge/discharge curves of WCM-Al2O3-Na2/3Ni1/3Mn2/3O2, and (c) cycle performance in a voltage range of 2.5−4.3 V at 0.5 C. (d) Rate performance, (e) TEM image of after-cycling Na2/3Ni1/3Mn2/3O2, and (f) TEM image of after-cycling WCMAl2O3-Na2/3Ni1/3Mn2/3O2. Reprinted with permission from ref 82. Copyright 2016 Elsevier.

J. H. Jo et al.83 modified the surface of Na2/3Ni1/3Mn2/3O2 with NaPO3. A NaPO3 coating layer is formed by the reaction of NH4H2PO4 and surface sodium residues of Na2CO3 and NaOH during melt-impregnation at 300 °C. As can be seen from Figure 15a−d, NaPO3-coated Na2/3Ni1/3Mn2/3O2 shows notably enhanced cyclability without alteration in the electrochemical reaction behavior. In a half-cell test, the capacity retention is increased from ∼66% to ∼80% on going from bare to NaPO3-coated electrodes. In a full-cell paired with hard carbon, the capacity retention for 300 cycles is ∼22% for the bare electrodes and ∼73% for the coated electrodes. The function of the coating layer can be summarized in two aspects. On one hand, the high ionic conductivity of NaPO3 could act as a fast diffusion channel for ion transport inside the grain boundary of the cathode84 as well as in the interface between the cathode and electrolyte, thus eliminating phase separation in the coated sample (Figure 15e,h). On the other hand, both structure and morphology are well retained with NaPO3 coating, as compared to the serious damage of particles observed in bare electrode (Figure 15f,g,i,k).

A thin layer of MgO was coated on the surface of Na0.5Ni0.26Cu0.07Mn0.67O2 by melt-impregnation to improve the interface property.85 However, the improvement in cycle stability was puny when the cell was cycled over the voltage range of 2.0−4.5 V at 0.33 C (45 mA/g). Based on these reports, we believe surface coating can lead to a slight improvement in the electrochemical performance of Na2/3Ni1/3Mn2/3O2 via suppressing the interfacial reaction, while doping can effectively improve the electrochemical performance via suppressing the drastic phase transitions of Na2/3Ni1/3Mn2/3O2 upon cycling, which is the most critical factor that causes its capacity to fade. Surface doping is expected to render the synergy of these two effects; thus, it should be a promising approach and is deserving of intensive investigation in the future. 3.4. Selection of Binder. Y. Yoda et al.32 adopted watersoluble poly-γ-glutamate (PGluNa) binder for Na2/3Ni1/3Mn2/3O2 electrode, taking advantage of the waterresistance property of Na2/3Ni1/3Mn2/3O2. The PGluNa electrode shows better electrochemical performance regarding capacity retention and Coulombic efficiency as well as rate 22061


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Figure 15. (a) Charge/discharge curves and (b) cycle performance of bare Na2/3Ni1/3Mn2/3O2 and NaPO3-coated Na2/3Ni1/3Mn2/3O2 over 1.5− 4.3 V at 0.1 C (20 mA/g). (c) Long-term cyclability and Coulombic efficiency of full cells assembled with a hard carbon anode over 1.4−4.2 V at 0.2 C (40 mA/g). (d) Rate performance of full cells with a constant charging rate of 0.2 C. Rietveld refinements of XRD patterns and TEM images for post-cycled bare Na2/3Ni1/3Mn2/3O2 (e, f, g) and NaPO3-coated Na2/3Ni1/3Mn2/3O2 (h, i, k). Reprinted with permission from ref 83. Copyright 2018 John Wiley and Sons.

methods.35,42,86,87 In the case of the sol−gel method, the crystal structure of the final product is closely related to the annealing temperature.35 P3-type structure tends to form at 850/800 °C, and a true P3 phase with NiO impurity is obtained at a further decreased temperature of 700 °C. Apart from the synthesis temperature, heating rate is another critical parameter affecting the final phase. A heating rate of 2 °C/min and an annealing temperature of 900 or 950 °C could lead to the pure P2 phase, while NiO impurity is clearly observed with a higher heating rate at same annealing temperature.35 We suppose this influence is also applicable to other synthesis methods. 4.2. Possible Oxygen Redox Reaction. D. H. Lee et al.42 first proposed the possible oxygen redox reaction involved in Na2/3Ni1/3Mn2/3O2 by calculating the density of states (DOS). Later, S. Doubaji et al.35 proposed that oxygen redox reaction takes place in P2-NaxCo1/2Mn1/3Ni1/6O2 based on the in situ XANES spectra. During charging from 4.2 to 4.5 V, the only minor shift in the Co K-edge spectrum cannot account for the large capacity of ∼80 mAh/g in this region. In addition, no

capability. In particular, the electrode using PGluNa binder displays initial Coulombic efficiency and capacity retention for 50 cycles of ∼95% and ∼89%, respectively, which is much better than the values of ∼80% and ∼71% for the electrode using poly(vinylidene difluoride) (PVdF) binder. The discharge capacity of PGluNa is ∼30 mAh/g higher than that of the PVdF electrode at 2 C. On the basis of XPS analyses, PGluNa binder serves to stabilize the electrode surface and suppress electrolyte decomposition during charging/discharging. Another reason for the better performance of PGluNa electrode is the reduced electrode resistance during cycling. The origins of lower resistance include intensive adhesion strength and moderate electrolyte penetration of the PGluNa binder, both of which lead to suppressed electric isolation of active particles.

4. NOTES AND CHALLENGES 4.1. Synthesis of Na2/3Ni1/3Mn2/3O2. P2Na2/3Ni1/3Mn2/3O2 can be prepared in various ways, including solid-state, sol−gel, co-precipitation, and spray pyrolysis 22062


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induced by the removal of Na ions, thus causing the instability of the P2 structure. Introduction of guest cations into Na2/3Ni1/3Mn2/3O2 is able to improve the cycle stability by (1) suppressing the structural change at high voltage as more Na ions are allowed to stay in alkali layers and (2) inhibiting Na+/vacancy ordering as the presence of guest cations gives rise to more disordering in the structure and fewer structural and electronic processes upon Na ions’ de-intercalation/ intercalation. Surface coating could also enhance the cycle performance of Na2/3Ni1/3Mn2/3O2 to some extent due to the reduced side reactions. Low cost and high volumetric energy density are prime requirements for stationary load-leveling applications. In this regard, layered Na2/3Ni1/3Mn2/3O2 is ideally suited for stationary ESS due to its high reversible capacity and compact structural framework, as well as the cost-effective Na resources. Most Na-based layered oxides require dry conditions during sample handling because of their hygroscopic character, while Na2/3Ni1/3Mn2/3O2 is stable in ambient air, which could lower the environmental demands and reduce the manufacturing cost. It is generally acknowledged that NIBs always have lower energy density than their lithium counterparts due to the heavier weight and higher potential of sodium. However, the possible oxygen redox reactions involved in Na 2 / 3 Ni 1 / 3 Mn 2 / 3 O 2 , Na 2 / 3 Ni 1 / 3 ‑ x Mg x Mn 2 / 3 O 2 , and NaxCo1/2Mn1/3Ni1/6O2 provide important scientific direction for designing high-energy-density NIBs. As discussed in this Review, some improvements in the electrochemical performance of Na2/3Ni1/3Mn2/3O2 have already been achieved, but challenges remain on the way to its commercialization. First and foremost, much more effort should be made to improve the cycle stability of Na2/3Ni1/3Mn2/3O2 over a wide voltage range (i.e., 2.5−4.4 V) because the current performance cannot satisfy the application requirements. In addition, attention should be paid to improving the Coulombic efficiency per cycle, as it is closely associated with the full-cell performance where the Na reservoir is limited. An in-depth understanding on the structural evolution and possible oxygen redox reaction of Na2/3Ni1/3Mn2/3O2 as well as the critical role of dopants is required, as this knowledge could in turn serve as guide to improve the cycle life and energy density of the materials. Last but not least, in addition to the electrochemical properties, it is also necessary to pay attention to the air stability and thermal stability of modified Na2/3Ni1/3Mn2/3O2, as those factors will significantly affect the production cost of the material and battery safety.

obvious change is observed in the shape of transition metal Kedge spectral features, including pre-edge peak, shoulder peak, and main peak, meaning the local structure is preserved and phase transition is averted. Therefore, the origin of capacity at voltage above 4.2 V is most likely oxygen reaction. During discharging NaxCo1/2Mn1/3Ni1/6O2 to 2.0 V, all the transition metals exhibit stable K-edge XANES spectra between 4.5 and 4.2 V, which further supports the speculation that reversible oxygen redox reaction contributes the capacity in the highvoltage region. In fact, oxygen ions have been confirmed to take part in the electrochemical reaction of a similar composition of Na0.78Ni0.23Mn0.69O2.88 Charge/discharge measurements show that Na0.78Ni0.23Mn0.69O2 delivers a fortuitously high capacity upon initial charging to 4.5 V, along with an irreversible long voltage plateau, which is ∼60 mAh/g higher than the theoretical value based on the Ni2+/Ni4+ redox couple. Structural study indicates that a P2−O2 phase transformation does not take place in Na0.78Ni0.23Mn0.69O2, even at a high charge state of 4.5 V. EELS and sXAS results suggest a gradient distribution of TM oxidation states from the surface to the bulk of cycled particles, which is tightly associated with the formation of oxygen vacancies at the particle surface. An oxygen redox reaction could afford additional capacity beyond the traditional transition metal redox couples. Therefore, it is of great significance to comprehensively investigate this kind of reaction within layered oxides. The insights gained through the study could also provide valuable guidance for designing highenergy NIB cathode materials. 4.3. Air Stability. Most Na-based P2-type layered oxides are prone to react with airmainly the CO2 and H2O in air and form Na2CO3·H2O and a hydrated Na-deficient phase. As a result, the air-exposed material exhibits much higher charge/ discharge polarization and lower capacity than rigorously airprotected material.26,89 P2-Na2/3Ni1/3Mn2/3O2 is air stable, which is indeed an intrinsic advantage for mass production. The stability of Na2/3Ni1/3Mn2/3O2 in air may be related to the existence of superlattice ordering of Ni atoms within TM layers, which induces a very strong interlayer interaction and inhibits uptake of H2O. The presence of Mn3+ might also aid the intercalation of H2O.33 Although different approaches have been carried out to stabilize the structure and capacity of Na2/3Ni1/3Mn2/3O2 upon cycling, air stability of modified Na2/3Ni1/3Mn2/3O2 was rarely reported. The introduction of new cations may affect the hydroscopicity of Na2/3Ni1/3Mn2/3O2; for example, water can be intercalated into P2-Na2/3Co1/6Ni1/6Mn2/3O2.33 Therefore, we should pay attention to this issue in addition to the electrochemical performance in the future studies.

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected].

5. CONCLUSIONS AND PERSPECTIVES In this Review, we discuss the developments and challenges for Na2/3Ni1/3Mn2/3O2 material. Phase transitions at different Na+/vacancy ordering and structural transitions caused by oxygen gliding at high voltage (e.g., ∼4.2 V) are the two major factors hindering the long-term cycle life of Na2/3Ni1/3Mn2/3O2, while interfacial side reactions at high voltage also contribute to the capacity decay of Na2/3Ni1/3Mn2/3O2. The ionic property of Na and strong Na+−Na+ in-plane repulsions lead to the formation of ordered superstructures upon removal/insertion of Na ions. Oxygen layers glide at high voltage to reduce the increased electrostatic repulsive force between neighboring transition metal layers

ORCID

Jiaolong Zhang: 0000-0003-3128-2154 Wenhui Wang: 0000-0002-2449-619X Baohua Li: 0000-0001-5559-5767 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by National Natural Science Foundation of China (Grant No. 51872157), Shenzhen Technical Plan Project (Grant Nos. KQJSCX201622063


Review


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0226191136, JCYJ20170412170911187, and JCYJ20170817161753629), Guangdong Technical Plan Project (Grant No. 2015TX01N011), Local Innovative and Research Teams Project of Guangdong Pearl River Talents Program (Grant No. 2017BT01N111), and China Postdoctoral Science Foundation (Grant No. 2019M650663).

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