Roles of surface structure and chemistry on electrochemical processes in lithium-rich layered oxide cathodes

Nano Energy 30 (2016) 580–602

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Roles of surface structure and chemistry on electrochemical processes in lithium-rich layered oxide cathodes

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Weifeng Weia, , Libao Chena, Anqiang Pana, Douglas G. Iveyb a b

State Key Laboratory of Powder Metallurgy, Central South University, Changsha, Hunan 410083, People's Republic of China Department of Chemical & Materials Engineering, University of Alberta, Edmonton, Alberta, Canada T6G 1H9

A R T I C L E I N F O

A BS T RAC T

Keywords: Lithium-rich layered oxides Surface structure and chemistry Electrochemical processes Structural evolution Voltage/capacity fade Surface modification

Li-rich layered oxides (LLOs) are promising cathode candidates for next generation Li-ion batteries, as they exhibit a higher reversible capacity ( > 250 mA h/g), enhanced safety and much lower cost. However, LLOs generally suffer from high first cycle irreversible capacity (IRC) loss, poor rate capability, and a substantial voltage decay over prolonged cycling. These major challenges are closely dependent on the surface structure and chemistry of LLO cathodes and, thus, different surfaces induce different irreversible reactions resulting in various levels of battery performance. This review presents the current understanding, as well as recent highlights, on the roles and fundamentals of surface structure in LLO cathodes, from a materials science perspective, concerning surface structural disorder in pristine LLO (antisites, composition segregation and crystallographic facets), the roles of surface structures on redox processes (oxygen evolution, cation activation and reversible anion redox reactions), surface structural evolution during the first cycle and long-term electrochemical operation, and surface modification strategies to stabilize the surface structure and to mitigate the performance degradation of LLOs. However, some fundamental problems remain yet ambiguous, especially with regard to characterization and understanding of the surface structure and chemistry in relation to synthesis conditions and composition, and charge transfer and ionic transport of the interfacial processes within LLOs. In order to exploit the potential of LLO cathodes, a clear understanding of these fundamental questions are essential to optimize the synthesis parameters and material properties.

1. Introduction Lithium-ion batteries (LIBs), charateristic of high energy density and long cycle life, have been extensively applied as a power source in various portable electronic devices. LIBs have been the focus of significant research attention with the intention to advance their application in the hybrid electric vehicle (HEV) and plug-in electric vehicle (PEV) market and in grid-scale energy storage applications. The development of cathode materials with improved capacity, rate capability, cycle life, safety, and low cost is one of the most demanding fields in LIBs. Li-rich layered oxides (LLOs) are a class of cathode materials that have recently attracted significant attention as potential cathodes for next-generation LIBs. The main advantage of LLO materials are their high reversible capacities (250–300 mA h g−1), representing a substantial enhancement over commercially available LIB cathodes: Layered LiCoO2 (~140 mA h g−1), LiNi1/3Mn1/3Co1/3O2 (150–160 mA h g−1), spinel LiMn2O4 (~120 mA h g−1), and olivine LiFePO4 (~160 mA h g−1) [1–6]. Other beneficial attributes of these LLO materials are the lower cost and improved safety owing to a mass ⁎

use of Mn content. LLO research dates from the early 1990s, when Thackeray's group synthesized lithium-enriched Li1.09Mn0.91O2 and introduced the idea of xLi2MnO3(1–x)LiMnO2 [7]. The unique high capacities of LLOs have attracted significant attention and extensive investigations have been conducted to understand the electrochemistry and phase transformations during electrochemical cycling, and to enhance the electrochemical performance through introducing heteroatoms into the crystal structure or via surface modification. Since the crystal structure of LLOs can substantially influence their electrochemical behavior, it is essential to clarify the crystal structure of LLOs. The crystal structure of LLOs can be described in the following three ways: (1) nanoscaled intergrowth of trigonal LiTMO2 (TM=any transition metal) and monoclinic [8–13]; (2) a single monoclinic Li2MnO3 structure with several variations [14–17]; and (3) a trigonal LiTMO2 phase with the existence of a superstructure [14,18]. The structural controversy is ascribed to the identification of long-range symmetry determined by diffraction methods such as X-ray diffraction (XRD) and neutron diffraction (ND). LiTMO2 (trigonal, R-3m) and Li2MnO3 (monoclinic,

Corresponding author. E-mail address: [email protected] (W. Wei).

http://dx.doi.org/10.1016/j.nanoen.2016.10.066 Received 21 September 2016; Received in revised form 31 October 2016; Accepted 31 October 2016 Available online 02 November 2016 2211-2855/ © 2016 Elsevier Ltd. All rights reserved.

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Fig. 1. (a) Crystal structure of the LiTMO2 lattice, space group of R-3m (left) and the Li2MnO3 lattice, space group of C2/m (right). (Reprinted with permission [13]. Copyright 2013, Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim). (b) XRD pattern of LLOs with a formula of 0.5Li2MnO3·0.5LiNi0.5Mn0.5O2. (Reprinted with permission [19]. Copyright 2015 The Electrochemical Society). (c) Bragg-filtered HAADF-STEM image of 0.6Li2MnO3·0.4LiNi0.45Mn0.525Mg0.025O2 containing Li2MO3 phase (blue) and LiMO2 phase (green). (Reprinted with permission [20]. Copyright 2012 American Chemical Society). (d) HADDF-STEM image of Li1.2Mn0.567Ni0.166Co0.067O2 showing domains with LiTMO2 and Li2MO3-like structural units in Region II. (Reprinted with permission [13]. Copyright 2013 WILEY-VCH, Verlag GmbH & Co. KGaA.) (e) Bright-field TEM image of Li1.2Ni0.2Mn0.6O2 and (f) typical D-STEM diffraction patterns collected in the areas in the bright-field TEM image. (Reprinted with permission [21]. Copyright 2012 American Chemical Society.).

formula unit. Nevertheless, the excess Li utilization involves an extra redox couple in the LLOs, and the potentials of most TM4+/TM5+ redox processes are far beyond the electrochemical stability window of the electrolytes. Extensive studies have, therefore, been carried out to understand the distinct redox process of LLOs. Thus far, it is well established that the exceptionally high reversible capacity of LLOs is due to their distinct first charge plateau above 4.5 V vs. Li+/Li, which is related to the oxidation of oxygen ions from the lattice in LLOs during lithium deintercalation (Fig. 2a) [29,30]. The oxidation process induces partial oxygen evolution and surface structural rearrangement, leading to the gradual variation in the redox process upon extended cycling. The first demonstration of oxygen evolution in gas form was reported by Thackeray's group in 2006 using in-situ differential electrochemical mass spectroscopy (DEMS) [29]. The evolution of oxygen gas was detected at charge potentials above 4.5 V (Fig. 2b). Carbon-containing gas evolution was also observed simultaneously, which was initiated from side reactions with electrolytes [29]. The evolution of O2 and carbon-containing gaseous species has also been reported in other research work [31,32]. The irreversible oxygen evolved at the electrode surface stimulates the surface structural evolution accompanied by cation migration between the Li/TM layers [33,34]. This process leads to the rearrangement of the surface structure (Fig. 2c), but the bulk structure does not experience significant changes [35]. Meanwhile, the oxygen radicals may attack the ring-structure of carbonate-based solvents [36], forming a solid electrolyte interphase (SEI) via this irreversible reaction [37,38]. The oxygen radical may further react with the byproducts, such as carbon-containing gases and water, to form Li2CO3, the main product of the side reactions [32]. All the irreversible reactions result in the deterioration of the electrochemical performances of the LLOs, including the reversible capacity, rate capability, and cycle stability. LLOs undergo high irreversible capacity (IRC) loss in the first cycle, poor rate capability, and a substantial voltage decay over prolonged cycles. These major challenges are closely related to the irreversible reactions mentioned above, which are dependent on the surface species and, thus, LLO materials with different surface morphologies induce distinct irreversible reactions, resulting in various levels of battery performance. There are some reviews on the development of LLO

C2/m), as shown in Fig. 1a, share the same oxygen closed packed lattices (ccp) [13]. The different space group of Li2MnO3 results from the cation ordering between the Li and Mn ions in the transition metal (TM) layers of the lattice, allowing an complete intermixing of the two structures. In XRD patterns, all R-3m peaks overlap with C2/m peaks except for the so-called superstructure peaks that appear at 20–25° 2θ (Cu Kα), as shown in Fig. 1b [19]. These peaks can be indexed with respect to C2/m and are a result of cation ordering in the TM layers (e.g., LiMn6) accompanied by the appropriate stacking along the c-axis [5]. Consequently, all XRD peaks of LLOs can also be indexed to the C2/m structure. Recently, more direct observations of the local structure of pristine LLO materials was achieved using high-angle annular dark field scanning transmission electron microscopy (HAADF-STEM) [13,20] and STEM diffraction (D-STEM) [21], but there is still no clear agreement among researchers. Boulineau et al. observed separated domains of Li2MnO3 and LiMnO2 in pristine LLO materials using HAADF-STEM (Fig. 1c) [20]. The bright and dark contrast represents the Li/TM ordered and TM segregated regions, respectively. This finding supported the claim that the C2/m and R-3m structures coexist inside the material as a nanocomposite structure. A similar observation was reported by Zhou's group, as shown in Fig. 1d [13]. However, some research groups have discovered direct evidence of solid-solution phenomenon in LLOs. For example, various studies have discovered that LLOs obey Vegard's law since the lattice parameters change with composition in a linear relationship [18,21,22]. HAADF-STEM characterization has also been employed to display the formation of a single C2/m structure at an atomic scale, as shown in Fig. 1e and f [21]. Considering that the nanocomposite and solid solution domains approach the nanometer scale, the distinction between a solid solution and composite would be blurred. The local structure of LLOs will likely be ambiguous since the final structure is greatly affected by various factors, including the type and composition of the TM species, the Li/ TM ratio, the synthesis route and conditions, as well as the material characterization methods [23–28]. The high reversible capacity of LLOs, exceeding 250 mA h g−1 and surpassing theoretical values on the basis of the redox capabilities of the TM cations, implies a potential utilization of over one Li ion per 581

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Fig. 2. (a) First charge-discharge profiles of 0.3Li2MnO3·0.7LiNi0.5Mn0.5O2. (Reprinted with permission [30]. Copyright 2004 Elsevier.) (b) In-situ demonstration of oxygen gas evolution during electrochemical charging of Li1.2Ni0.2Mn0.6O2. (Reprinted with permission [29]. Copyright 2006 American Chemical Society.) (c) Atomic resolution HAADF-STEM image of electrochemically cycled Li[Ni1/5Li1/5Mn3/5]O2 in the surface region. (Reprinted with permission [35]. Copyright 2011 The Royal Society of Chemistry.).

2.1. Antisite defects

cathodes and the understanding of structure-electrochemical property relationship [2,5,19,39–44]; however, a critical review dedicated to the roles and fundamentals of surface structure in LLO materials, from a materials science perspective, is still lacking. In the present review, we will initially focus on surface structural disorder of the pristine material (antisites, composition segregation, and crystallographic facets), the roles of surface structures on redox processes (oxygen evolution, cation activation, and reversible anion redox reactions), surface structural evolution during the first cycle and long-term electrochemical operation, and surface modification strategies to stabilize the surface structure and to mitigate the performance degradation of LLOs. This will be followed by a summary and outlook regarding the remaining fundamental and practical challenges. Herein, we emphasize on the materials science aspects of surface structures in LLO cathode materials for LIBs. Some other essential issues such as electrolyte development, interfacial design, and device engineering are not included in this work.

Antisite defects are the most common form of disorder in layered cathode materials with two dimensional (2D) Li diffusion channels [46,47]. Ceder's group recently reported a LLO material that has a significant amount of cation intermixing, but displays superior electrochemical performance, which is associated with fast Li diffusion via hopping through intermediate tetrahedral sites [48,49]. It was also shown that in Li1.2Cr0.4Mn0.4O2 that cation intermixing makes no negative impact on the reversible capacity, but stabilizes the bulk structure [50]. With the assistant of aberration-corrected STEM and electron energy loss spectroscopy (EELS), Dixit's group shows that there exists a significant amount of antisite defects in Li1.2Ni0.175Mn0.525Co0.1O2, with Ni substituting on Li sites [51]. The antisite defects display a strong segregation tendency to the facets that are vertical to Li/TM layers rather than those are parallel to Li/TM layers, as shown in Fig. 3a–b [51]. This tendency is further confirmed by first-principle calculations [51]. More recently, Shukla's group systematically observed four pristine LLO materials with various morphologies and compositions. It has been shown that the primary particle surfaces are Co-, Ni-enriched spinel layers with some antisite defects that are crystallographic facet dependent (Fig. 3c–d) [17]. All these findings indicate that certain facets may be more effective to restrain surface reconstruction and to stabilize the structure of the LLO materials through manipulation of the exposed surface.

2. Surface structure of pristine LLOs Actual surfaces of pristine LLOs usually contain structural disorder, such as antisite defects, composition gradients, preferentially exposed crystallographic facets, etc. The types and concentration of surface structural disorder appearing in pristine LLOs are influenced by the chemical compositions, synthesis methods, and the applied synthesis parameters. These disordered surface regions may have a pronounced impact on the transfer and storage of electrons and ions in layered oxide cathodes [45]. Understanding the surface structure and controlling the surface disorder enable the regulation of ion transport, storage, and cycle performance of LLO cathodes.

2.2. Surface cation segregation Transition metal element enrichment and inhomogeneity are other kinds of surface disorder in layered oxide cathodes. It has been revealed that compositional variation in surface and bulk regions exists in pristine LLOs, which is facet sensitive [52,53]. Wang's group 582

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Fig. 3. HAADF-STEM images and EELS line scan profiles for a LLO particle along the (a) [110] and (b) [001] direction. Surface reconstruction and segregation of Ni/Co and antisite defects are observed in the reconstructed layers. (Reprinted with permission [51]. Copyright 2014 American Chemical Society.) (c) and (d) HAADF images and simulated STEM images of spinel surface and the layered bulk along [110] and [010] directions, respectively. (Reproduced with permission [17]. Copyright 2015 Macmillan Publishers Limited.).

electrochemical performance of the LLO materials may inspire the design of novel LLO materials with enhanced electrochemical performance.

proposed that the surface layer in Li1.2Ni0.2Mn0.6O2 nanoparticles shares the same structure as the bulk, but shows a different crystallographic orientation (Fig. 4a). More recently, they reported that Li1.2Ni0.2Mn0.6O2 prepared by hydrothermally assisted methods has slightly Ni-enriched surfaces and exhibits improved capacity retention and suppressed voltage decay during extended cycling, compared with materials produced by co-precipitation or sol–gel processes (Fig. 4b) [52,53]. TM element segregation to oxide surfaces is related to the nucleation and growth processes of TM oxide precursors during coprecipitation or sol–gel processes. More recently, Wang's group has provided direct evidence concerning the facet- and compositiondependence of surface segregation of Ni and Co [54]. Essentially, Ni segregates to (200)m facets to form a spinel-like structure, while Co preferentially segregates to (20-2)m facets and develops a rock salt-like structure; m refers to monoclinic C2/m symmetry, as shown in Fig. 4c [54]. Furthermore, segregation of Ni and Co is eliminated by raising the Ni content to produce Ni-rich LLOs [54]. The understanding on the relationship between the nanoscaled chemical distribution and the

2.3. Preferentially exposed crystallographic plane The exposed crystallographic planes on the cathode surface are critical to Li+ transport [55,56]. In LLOs with a layered structure, only the {010} planes offer accessible diffusion paths for Li+ cations. Sun's group has revealed that enhanced rate capability was achieved in LLO cathode materials through raising the amount of {010} exposed facets [57,58]. Note that it remains challenging to manufacture LLOs with preferentially exposed {010} planes because these planes are highenergy facets and tend to vanish during crystal growth process [57]. More recently, in order to take the advantages of the hierarchical architecture and the controlled active exposed planes of the primary particles, Wu's group demonstrated a rational approach for fabricating a hierarchically structured Li1.2Mn0.6Ni0.2O2 cathode with high per583

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Fig. 4. (a) Three dimensional (3D) energy dispersive X-ray (EDX) mapping and atomic resolution HAADF-STEM image of LiNi0.5Mn0.5O2 nanoparticle. (Reprinted with permission [52]. Copyright 2012 American Chemical Society.) (b) EDX mapping and electrochemical performance of Li1.2Ni0.2Mn0.6O2 prepared by the hydrothermally assisted, co-precipitation and sol–gel processes. (Reproduced with permission [53]. Copyright 2014 American Chemical Society.) (c) EDX mapping and atomic resolution HAADF-STEM images to show the plane selectivity of Ni and Co SSLs. (Reprinted with permission [54]. Copyright 2016 WILEY-VCH, Verlag GmbH & Co. KGaA.).

3. Role of surface structure on electrochemistry

centage of {010} exposed planes in the surface areas. This hierarchical cathode material exhibited excellent rate performance and remarkable capacity retention due to the enhanced Li diffusion rate [59].

One of the critical characteristics of LLOs is the distinct chargedischarge profile of the initial cycle, as shown in Fig. 2a. During the sloping part in the first charge profile, Li+ extraction is associated with the oxidation process of the TM cations to the 4+ chemical state. The plateau part appears above 4.5 V and consists of Li being extracted concurrently with oxygen, which can be viewed as losing Li2O. The discharge profile of the first cycle is rather complicated owing to the unique first charge process. During the initial discharge process, the oxidized TM cations in the 4+ oxidation state are restored to their original chemical states. To compensate for the O2− loss in the first charge process, alternative reduction process is required; some possible mechanisms have, therefore, been proposed to account for the exceptionally high capacities of LLOs ( > 250 mAh g−1). The most widely accepted one is that some Mn4+ cations are reduced to Mn3+ cations. Some other working mechanisms, including reversible oxygen redox and side reactions with the electrolytes, may deliver additional discharge capacity [60–65].

2.4. Summary on the surface structures formed in pristine LLO materials Based on recent research findings, it is apparent that the complex surface strucutures of pristine LLO materials are rather sensitive to the composition, synthesis routes and employed synthesis parameters. It would be of interest to determine how the surface structure evolves if the content of Li and Mn is reduced and/or annealing temperature and time are lowered, which provides some new directions for material design and synthesis routes in this class of materials. In addition, with a comprehensive appreciation of the surface structures, including the chemistry and their crystallographic correlation with the bulk material, the researchers may attempt to explore the effect of different surface structures on initial delithiation/lithiation process and subsequent extended cycling, and to provide some strategies that probably resolve the technical aspects such as fast voltage decay and capacity fade and poor rate capability. 584

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Fig. 5. Schematic illustrations of the proposed reaction mechanisms in the Li1.2Co0.13Ni0.13Mn0.54O2 electrode surface. (Reprinted with permission [66]. Copyright 2011 American Chemical Society.).

Fig. 6. (a) Differential electrochemical mass spectroscopy analysis detecting O2, CO and CO2 gases during the initial three electrochemical cycles of Li1.2Ni0.2Mn0.6O2. (Reprinted with permission [32]. Copyright 2012 American Chemical Society.) (b) Optical microscopy images of the surface of the pristine Li1.2Ni0.13Co0.13Mn0.54O2 electrodes (left, top), discharged to 3.0 V after 4.8 V charge (right, top), and fully discharged to 2.0 V after 4.8 V charge (bottom). (c) XPS O 1s (left) and C 1s (right) local scan of the surface of a Li1.2Ni0.13Co0.13Mn0.54O2 electrode in the pristine state and at 4.8 V-charge, 3.0 V-discharge, 2.0 V-discharge, following 4.0 V-charge, and 4.8 V-charge (from bottom to top). (Reprinted with permission [66]. Copyright 2011 American Chemical Society.) (d) FT-IR spectra of Li1.2Ni0.2Mn0.6O2 electrodes at various states of charge, (Reproduced with permission [32]. Copyright 2012 American Chemical Society.) and (e) Raman mapping of Li1.2Ni0.2Mn0.6O2 after discharging to 2.0 V, showing 1090 cm−3 Li2CO3 peak. (Reprinted with permission [71]. Copyright 2012 American Chemical Society.).

established TM redox reactions, therefore, the electrochemical redox process of the oxygen species occurred at the surface contributes a portion of the reversible capacity of the LLOs. More recently, Delmas's group has suggested a working mechanism to explain the reversible oxygen ion oxidation in the bulk and irreversible oxidation (oxygen loss) at the surface [60,61,67,68]. On the basis of electrochemical, structural, and spectroscopic measurements, a densification process was confirmed that the oxygen evolution at the surface and the migration of TM cations from the surface to the bulk occur simultaneously, but no significant changes are observed in the bulk structure [60,61,67,68]. Since the oxygen gas participates in electrochemical/chemical reactions with electrolytes, oxygen evolution from the LLO surface plays a critical role in closed electrochemical cells [69–71]. Kang's group reported that a significant amount of oxygen as well as carbonbased gaseous species are evolved in the charge-discharge process of

3.1. Oxygen evolution from the LLO surface In most LLOs, oxygen gas is generated from the oxide surface during the voltage plateau of the initial charge process. It is generally recognized that high discharge capacity of the LLOs after the first charge is due to two major mechanisms: one is the activated manganese redox process (Mn3+/Mn4+) and the other is the oxygen reduction reaction occurred at the surface. Komaba's group summarized the proposed reaction schemes, as shown in Fig. 5 [66]. During the first charge process, a considerable volume of oxygen vacancies is generated at the surface due to oxygen evolution process. Consequently, the coordinated TM cations migrate to the octahedral vacancies in Li layers, resulting in the formation of a surface reconstructed layer. During the subsequent discharge process, in particular below 3.0 V, the generated O2 molecules could be electrochemically reduced to O2− anions, accounting for the extra capacity [66]. Apart from the 585

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charge or generated from the electroreduction of oxygen gas molecules at the electrode surface during discharge. The ring-structured carbonate-based solvents are known to be vulnerable to chemical attack by oxygen radicals, breaking the ring structure [36]. Therefore, various byproducts, such as carbon dioxide, carbon monoxide, and water, are formed. The oxygen radical can further react with the byproducts, which finally form Li2CO3, the main product of the side reactions (Fig. 6d and e) [71]. 3.2. Transition metal cation activation originating from the surface The extra capacity in the first charge process can be ascribed to the concurrent extraction of Li+ and O2−, whereas the high reversible capacity for subsequent cycling can be explained by the activation of additional redox couples in the LLOs. The high reversible capacity of LLOs is associated with the low-voltage range of 3.0–3.2 V after charging over ∼4.5 V vs. Li/Li+, as shown in cyclic voltammograms or differential capacity vs. potential (dq/dV) plots (Fig. 7a and b) [66,72]. The redox process at lower potential is attributed to the Mn3+/ Mn4+ redox couple that is inactive in the pristine state, which is illustrated with the band structures for the LLOs in Fig. 8 [19]. Numerous studies have focused on the identification of the activation of the Mn3+/Mn4+ redox reactions using various techniques. Komaba's group showed that the X-ray absorption near edge structure (XANES) of the Mn K-edge shifted to lower energy after electrochemical cycling, indicative of the reduction of Mn4+ to Mn3+(Fig. 9a) [66]. In contrast, the XANES of the Ni K-edge recovered almost reversibly to the initial state (Fig. 9b) [66]. A similar tendency was also observed by other research groups [73,74], as shown in Fig. 9c and d. It is worthy to note, however, that X-ray absorption methods provide primarily average coordination chemical information and they have inherent limitations on local structure and chemical information determination. STEM, as an alternative approach, is very appropriate for local area microanalysis. Recently, Meng's group has shown, through EELS in an aberration-corrected STEM, that the Mn cations near the surface (~2– 5 nm) were reduced to a lower valence state and the oxygen K-edge pre-peak signal diminished after electrochemical cycling (Fig. 10a and b), suggesting that the activation of the Mn3+/Mn4+ redox couple originates from the LLO surface [35]. In addition, they also demonstrated the presence of Mn3+ in the Mn L-edge using EELS and in the Mn 2p spectra of the discharged LLC using X-ray photoelectron spectroscopy (XPS), as presented in Fig. 10c and d [75]. In the discharge process, i.e., the re-insertion of the Li ions, it is considered that owing to the lack of electron compensation of Ni cations in the LLOs, the Mn cations are activated as electrons filled in the vacant electronic bands above the Ni3+/Ni4+ and Ni2+/Ni3+ bands, i.e., the

Fig. 7. (a) Cyclic voltammogram of 0.3Li2MnO3·0.7Li(MnCoNi)1/3O2. (Reproduced with permission [72]. Copyright 2008 American Chemical Society.) (b) Differential capacity vs. voltage (dq/dV) plots of Li1.2Ni0.133Mn0.533Co0.133O2. (Reproduced with permission [66]. Copyright 2011 American Chemical Society.).

LLOs, as shown in Fig. 6a [32]. These suggested electrochemical/ chemical reactions involving oxygen radicals continuously produce lithium carbonate during the charge-discharge process. Komaba's group also demonstrated the formation of byproducts at the electrode surface after discharge of LLOs below 3.0 V, using optical microscopy, and revealed that the byproducts include carbonate (CO32−) as well as C-H and CH2-components (Fig. 6b and c) [66]. The formation of carbon-based byproducts results from the decomposition of the carbonate-based electrolyte solvents by oxygen radicals [32]. The oxygen radicals are mostly extracted from the lattice of LLOs during

Fig. 8. Schematic electronic density of states (DOS) of LLO cathode at different states of charge. LLO undergoes oxygen evolution and structural evolution accompanied with electronic structure rearrangement resulting in the Mn3+/Mn4+ redox activation. (Reproduced with permission [19]. Copyright 2015 The Electrochemical Society).

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Fig. 9. The activation of the Mn3+/4+ redox reactions during extended cycling of LLOs. Ex-situ XANES spectrum of (a) Mn K-edge and (b) Ni K-edge of Li1.130Ni0.203Co0.203Mn0.464O2. (Reproduced with permission [66]. Copyright 2011 American Chemical Society.) Half-height energy of (c) Mn K-edge and (d) Ni K-edge of in-situ XANES analysis of Li1.2Ni0.17Co0.07Mn0.56O2. (Reproduced with permission [73]. Copyright 2011 Elsevier.).

Mn3+/Mn4+ band (Fig. 8) [19].

also detected that the reductive coupling mechanism with stabilized oxygen anions was highly favored in the presence of Sn4+ as compared with other 3d TMs such as Mn [62].

3.3. Reversible anion redox reactions In addition to the TM activation process, a new redox chemistry has been proposed to explain the high rechargeable capacity for LLOs [60,61,76]. Delmas's group proposed that a reversible oxygen redox reaction occurs in the bulk region of the large sized particles of Li1.2Mn0.54Co0.13Ni0.13O2, while irreversible oxygen loss occurs in the surface region as indicated by redox titration, stoichiometric analysis, and structural analysis [60,61]. Oishi's group further demonstrated the presence of the oxidized oxygen species in the charged materials and their reversible reduction in the discharged states by using soft X-ray absorption spectroscopy (XAS) [76]. Recently, it was proposed for LLOs containing 4d metal elements, such as Ru and Sn, that the reversible reaction of oxygen species (2O2−/(O2)n−) occurs in the first charge/discharge process instead of irreversible O2 evolution, when the orbitals of TM ions (3d or 4d) and neighboring oxygen ions (2p) strongly hybridize each other [62,63]. Tarascon's group observed the oxidized form of oxygen species in charged Ru-containing LLOs using XPS (Fig. 11a) and electron paramagnetic resonance (EPR) (Fig. 11b) [63]. The XPS O 1s spectrum of Ru-containing LLOs charged to 4.6 V revealed the existence of oxidized lattice oxygen (O22−). After the first discharge process, the XPS O 1s spectrum restores to the original one, suggesting the disapperance of O22− [62]. The EPR spectra at room temperature (Fig. 11b, top) show a narrow isotropic line at 4.0 V, implying the existence of Ru5+. At 4.6 V, the line shape widens asymmetrically, indicative of the formation of additional unpaired spin sources, i.e., charged oxygen species such as O23− or O2− states. The peak disappeared when discharged to 2.0 V, suggesting the reversibility of the redox activity of oxygen. The EPR spectra collected at 4 K further exhibit the difference between the paramagnetic species at 4.0 V and 4.6 V charge (Fig. 11b, bottom), providing evidence of the existence of peroxo/superoxo-like species with a strong distortion of the MO6 octahedra [77]. Tarascon's group

3.4. Summary on surface-related reactions and mechanisms One of the most distinct features of the LLOs is the oxygen evolution process that occurs at the first charge more than 4.6 V. This oxygen evolution process comprises a reversible redox process of 2O2−/(O2)n− and an irreversible process of lattice O2− loss [32,62]. The high specific capacity of the LLOs is ascribed to the reversible process, whereas the irreversible loss of lattice O2− results in a variety of unwanted chemical reactions at the surface and in turn, continuous structural evolution on extended cycles. The origin of the involvement and reversibility of the oxygen-containing species is still inconclusive. Moreover, even though the unconventional anion participation and/or transition metal cation activation lead to higher reversible capacities, several related challenges that hinder electrochemical kinetics and cycling stability remain in the LLO materials. Owing to the high charge voltage, the side products formed on the electrode surface through the complex reactions occur between oxygenated species and the organic electrolyte. An unambiguous understanding of the working mechanism of complex interfacial reactions that occurs along the LLO surface, together with quantitative evaluation of their contributions to the irreversible and reversible capacities will definitely inspire the advance of the LLO cathode materials. Moreover, how to suppress the irreversible oxygen evolution and to protect the electrode surfaces while promoting the anion redox process will be critical to promote the electrochemical cyclability of the LLO materials. Nevertheless, in spite of the importance of surface-related reactions, the phase transformations occurred to the surface and bulk materials are equally important to the electrochemical performance of the LLO materials. 587

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Fig. 10. (a–b) EELS spectra of the O K-edge and Mn L-edge of pristine and cycled Li(Ni1/5Li1/5Mn3/5)O2 in surface region. (Reprinted with permission [35]. Copyright 2011 The Royal Society of Chemistry.) (c–d) EELS and XPS spectra taken from the Li(Ni1/5Li1/5Mn3/5)O2 at different discharge states. (Reprinted with permission [75]. Copyright 2013 The Royal Society of Chemistry.).

0.5Li2MnO3-0.5LiNi1/3Co1/3Mn1/3O2 after the first charge process to 4.8 V [66]. In addition, significant broadening and reduced intensity was observed in the XRD reflections corresponding to the superstructural Li/Mn cation ordering in the TM layers after the first charge (Fig. 12a) [60]. This indicates irreversible structure reconstruction of the LLO electrodes along with the Li extraction from the TM layers. It was proposed that oxygen evolution occurs at the surface accompanied by the removal of Li in the TM layers, which induces TM cation migration into Li vacant sites from the surface region to the bulk, resulting in lattice densification to a normal layered structure, i.e., a LiMO2-like structure [33]. It was estimated that ∼5% of the TM cations simultaneously migrated from the TM layers to the Li layers after oxygen evolution in the Li2MnO3-based LLO system [66]. It is worthy noting, however, that the TM ordering in the bulk after the first cycle remains almost the same as that of the pristine material based on HADDF imaging along the [100] zone axis (Fig. 12b) [84]. More recently, other research work has been performed to probe structural evolution on the oxide surface after the first charge [50,84]. Boulineau's group confirmed, based on HAADF-STEM and EELS analyses, the presence of an extra phase with a defect spinel structure on the Li1.2Mn0.61Ni0.18Mg0.01O2 surface after the first cycle (Fig. 12c) [84]. A layered-defective spinel structural transformation occurs gradually from the surface to the bulk of the LLO particles. Actually, TM cations were observed to be in the inter-slab space at the surface region, whereas they are only positioned in the TM slabs in the bulk region [84]. In Cr-based LLO materials, after the first charge process, there is a significant variation between the surface and the bulk

4. Surface structure evolution during charge-discharge process LLOs experience a substantial change in the electrochemical profiles after the distinct first charge process [20,34,66]. It has been demonstrated that oxygen oxidation occurs in the high-voltage charge region, leading to irreversible oxygen gas evolution [29,66]. The oxygen redox reactions induce a change in the surface structures and the Li chemical potential, so the subsequent electrochemical profiles change after the first charge process [66,78,79]. Moreover, the charge-discharge profiles undergo steady changes during long-term cycling, with a gradual decay of the average voltage [80–83]. The voltage decay is attributed to irreversible structural evolution from layered to spinel- or rock salt-type phases initiated from the LLO surface. In this section, therefore, we review the previous reports concerning the surface structural transition of LLOs during the first cycle and long-term cycling process and the associated degradation in electrochemical performance, including large first cycle IRC loss, voltage decay, and poor rate capability. 4.1. Surface structural reconstruction after first charge process: the origin of first cycle IRC loss The irreversible evolution of oxygen gas during the first charge of LLOs brings about the formation of oxygen vacancies in the LLO lattice and the migration of TM cations. It was revealed by XRD Rietveld analysis that ∼7.5% of the oxygen vacancies exist in the lattice of 588

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Fig. 11. (a) The deconvoluted XPS O1s spectra of Li2Ru0.5Sn0.5O3 in the state after charging to 4 V and 4.6 V, and after discharging to 2 V (from top to bottom) and (b) X-band EPR spectra taken at room temperature (top) and 4 K (bottom). (Reprinted with permission [63]. Copyright 2013 Nature Publishing Group.).

Fig. 12. (a) XRD patterns of pristine and charged/discharged LLO materials. (Reprinted with permission [60]. Copyright 2013 The Electrochemical Society.) (b) Atomic resolution HADDF-STEM image in [100] zone axis taken from the bulk region of LLO material after 1st cycle. (c) Atomic resolution STEM-HAADF image, EELS mapping, and Mn/Ni concentration taken from the surface part of LLO material after 1st cycle. (Reprinted with permission [84]. Copyright 2013 American Chemical Society.).

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Fig. 13. (a) HAADF- and (b) ABF-STEM images of LLOs at the 2.0 V discharge state. The corresponding line intensity profiles for the TM/Li layer, oxygen layer, and lithium layer from the HAADF (c) and ABF images (d). (e) and (f) Line profiles for lines 7−16 in the surface areas of the HAADF and ABF images. (Reprinted with permission [50]. Copyright 2015 American Chemical Society.).

Fig. 14. Atomic resolution HAADF-STEM images taken from pristine and cycled LLO particles. (a) Identical structure from surface to bulk in the pristine sample. (b) and (c) Gradual increase of the thickness of the surface reconstructed layer. (d) The spinel phase speads from the surface to the core region. (e) Atomic resolution STEM-HAADF image and corresponding fast Fourier transformation (FFT) pattern along the [101] direction. (f) Atomic resolution STEM-HAADF image to show the coexistence of I41 structure and spinel structure in a cycled material. (Reprinted with permission [87]. Copyright 2015 American Chemical Society.).

detected in Li layers, while in the surface area, obvious dots are visible in the Li layers, suggesting that a portion of the TM cations have migrated into the Li layers. These structural rearrangements and oxygen evolution closely correlate with Mn4+/Mn3+ activation during discharge due to reconstruction of the electronic structure.

regions. The two regions are divided by the blue dashed line in the HAADF- and annular bright field (ABF)-STEM images, as shown in Fig. 13a and b. The corresponding line intensity profiles of the Crbased LLO at the 2.0 V state during discharge are shown in Fig. 13c–f [50]. In the bulk region, no peaks corresponding to the TM cations are 590

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Fig. 15. HADDF-STEM images of 0.4Li2MnO3·0.6LiNi1/3Co1/3Mn1/3O2 after 100 cycles along the [310] monoclinic zone axis. (Reprinted with permission [88]. Copyright 2014 WILEYVCH, Verlag GmbH & Co. KGaA.).

Within the oxygen-loss plateau in the first charge, the removal of Li2O from the lattice generates Li and oxygen vacancies in the Li/TM layers [33,85]. In the subsequent discharge process, a certain amount of the generated oxygen vacancies and Li vacancies are removed due to TM migration to the vacant sites, i. e., the structural densification process [86]. However, owing to the irreversible process, not all the deintercalated Li cations in the initial charge process can be reintercalated into the LLO lattice, leading to a large first cycle IRC loss. The residual oxygen vacancies would be eliminated through TM cation migration upon extended electrochemical cycles, which might demonstrate as steady capacity fading in LLO materials, i.e., IRC loss [86].

are randomly intermixed. The atomic arrangement of the rock salt phase was observed on the particle surface in cycled LLOs using HADDF-STEM [88]. It is apparent that the rock-salt phase consists of more compact TM ordering when compared with the spinel phase, as shown in Fig. 15 [88]. Furthermore, the rock salt phase is mainly found in the surface region and coexists with the spinel-like phase in the bulk region. Thus, Cho's group proposed that the observed spinel-like structure may be an intermediate phase between the layered and cubic rock salt structures in the cycled LLOs [88]. A similar phase transformation sequence from a C2/m to R-3m structure after initial activation, then to an Fd-3m spinel-like structure and finally to an Fm-3m rock salt structure, was also revealed by Wang's group [89]. The formation of the rock salt phase during cycling is detrimental to the capacity retention due to the absence of a Li diffusion path in the lattice [90]. Thus, the surface structural changes from layered to spinel or rock salt phases in cycled LLOs has been regarded as the origin of the gradual electrochemical degradation, such as voltage decay and low rate capability. The structural transformation from the layered to spinel or rock salt phases is closely associated with TM cation migration between the TM and Li layers, which involves TM cations passing through the adjacent tetrahedral sites [91]. The magnitude of structural transformation depends on the type and oxidation states of the TM cations due to the different tetrahedral site energies of the transition states [46,92]. As the oxidation state of the TM cations increases, a stronger repulsion exists between the TM cations in the tetrahedral sites and neighboring octahedral sites [92]. Based on detailed HAADF-STEM analysis, Tarascon's group compared the number of TM cations trapped in interstitial tetrahedral sites in Li2Ru0.75M0.25O3 systems (M=Ti, Ru, or Sn) and revealed a correlation between the trapping amount and ionic size of the TM, as shown in Fig. 16 [93]. They concluded that the degree of cation trapping in tetrahedral sites increases with a smaller ionic size, resulting in the faster voltage decay of Li2Ru0.75Ti0.25O3 compared with the other materials containing cations with larger ionic sizes, such as Ru, or Sn [93]. The surface structural transition for LLOs, originating from TM interlayer migration, is considered as the main reason for the voltage decay, poor rate capability, and capacity fade during extended electro-

4.2. Correlation between surface structural evolution and voltage/ capacity decay Upon continuous electrochemical cycling, atomic-scale structure analyses have provided direct evidence for the structural transformation on the surface of LLO cathodes [8,20,35]. In Fig. 2d, the HADDFSTEM image of electrochemically cycled Li1.2Ni0.2Mn0.6O2 reveals an evident difference in the [001] atomic contrast on the surface [35]. This finding represents the structural transformation from the layered- to spinel-type structure due to the migration of TM ions into the Li layer. More recently, Wang's group proposed a more detailed surface structural transformation that involves progressive enrichment of TM and Li depletion upon electrochemical cycling (Fig. 14) [87]. It is therefore considered that the structure change of the surface reconstructed layer with an I41 structure is driven by its composition gradient [87]. Moreover, with more Li cations extracted from the LLO lattice, a M3O4-spinel was grown along the I41 structure (Fig. 14f), suggesting a I41-spinel structural transformation. Thus, Wang's group proposed a structural transformation sequence of C2/ m→I41→M3O4-spinel for the LLO surface layer upon extended cycles [87]. It is worth noting that the spinel phase is not constrained to the surface region and may spread to the bulk after extended cycling. In addition to the spinel phase, a cubic rock salt structure was identified in the LLOs after extended electrochemical cycling. The rock salt structure is considered as a disordered form of the layered structure between the TM and Li layers, i.e., all the Li and TM sites 591

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Fig. 16. Atomic resolution HAADF-STEM images taken from the cycled Li2Ru0.75M0.25O3 (M=Sn or Ti) along the [1–10] direction. The white arrows and green contrast profiles, and yellow contrast profiles show TM cations trapped in the tetrahedral and octahedral sites in the Li layers, respectively. (Reproduced with permission [93]. Copyright 2015 Nature Publishing Group.).

perative to develop an understanding of the nuance and complicacy in a wide length scale, and in turn, to develop methodologies to alleviate the structural changes occurred to the surface and subsurface regions. For instance, in addition to the crystallographic information, the chemical and morphological information of the subsurface phases would also be equally important. It is recognized that the surface changes may come from bulk chemistry gradient that spreads to the surface and the interfacial reactions occurred along the surface/electrolyte interfaces. It will be essential to differentiate and quantify their contribution to surface changes to develop effective strategies to reduce the surface structural evolution of LLOs.

chemical cycling. Firstly, gradual migration of TM cations especially Ni cations from the bulk to the surface will bring about depletion of Ni in the bulk structure and in turn, leads to capacity fading. Secondly, the surface becomes TM-rich since the TM cations gradually occupy the Li position in the Li slabs. Based on the simulation study by Ceder's group [48], the energy barrier for Li migration will increase accordingly, giving rise to the voltage fade. Thirdly, the formation of spinel or rock salt structures at the surface may effectively reduce the Li transport path, leading to both voltage and capacity fading. Finally, the dissolution of the TM cations, especially along certain crystal facets, leads to corrosion of LLO particles (such as the formation of cracks and pits) and, consequently, capacity fading. Therefore, it is crucial to optimize the surface of LLOs to obtain cathode materials with high reversible capacity, voltage stability, and rate performance.

5. Modification of surface structure and chemistry Considerable technical challenges are associated with LLOs, including the above-mentioned large IRC loss and voltage decay, as well as poor rate performance, which are associated with the surface structural transition on LLOs originating from transition metal interlayer migration. Therefore, surface modifications are considered to be major strategies to improve electrochmical cyclability, rate performance, and even thermal stability of LLOs. On the basis of proposed mechanisms to explain the main positive impacts on the electrochemical performance of LLO materials, surface modification approaches can be categorized as follows: (1) controlled growth of preferential orientation active planes that enhances Li+ diffusion on the oxide

4.3. Summary on surface structure evolution upon electrochemical cycling The surface layer reconstruction after first charge process and gradual structural transformation from the surface to subsurface regions are generally considered irreversible, which account for the origin of first cycle IRC loss and steady voltage/capacity fade of the LLO materials, respectively. However, diverse results on structural evolution process were observed in the LLO materials with various chemistries and microstructures. Sustained efforts are therefore im592

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Fig. 17. (a) Schematic illustrations of two types of nanoplates with (001) and (010) exposed facets. (b) Typical morphology of the crystal habit-tuned nanoplate LLO material. (Reprinted with permission [57]. Copyright 2010 WILEY-VCH, Verlag GmbH & Co. KGaA.) (c) and (d) Hierarchical Li1.2Mn0.6Ni0.2O2 cathode consisting of nanoplates with a high surface area of exposed {010} planes. (Reproduced with permission [59]. Copyright 2014 WILEY-VCH, Verlag GmbH & Co. KGaA.) (e) and (f) SEM image and XRD pattern of hierarchical Li1.2Mn0.54Ni0.13Co0.13O2 spheres. (Reproduced with permission [99]. Copyright 2014 American Chemical Society.).

electrolytes; (5) multi-functional electrochemically active layer that protects the LLO surface from HF corrosion and stabilizes the bulk structure; and (6) gradient surface doping to stabilize the bulk structure and protect the LLO surface.

surface; (2) electrically conducting coatings that facilitate charge transfer on the oxide surface; (3) fast Li+-conducting layer that accelerates Li+ diffusion on the oxide surface; (4) electrochemically inactive protective layer to provide both HF scavenging and physical protection to reduces transition metal dissolution from the oxide surface and to suppress side reactions between LLOs and non-aqueous

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Fig. 18. (a) Typical morphology of LLO samples with a rGO/AlPO4 coating. (b) Enhanced rate performance of rGO/AlPO4-coated LLOs. (Reprinted with permission [108]. Copyright 2014, The Royal Society of Chemistry.) (c) and (d) Typical high-resolution transmission electron miscroscopy (HRTEM) image and schematic view of the Layered@Spinel@Carbon heterostructure in LLOs. (Reprinted with permission [109]. Copyright 2015, The Royal Society of Chemistry.).

surface energy of the crystallographic surfaces [100]. They proposed that this fusiform LLO material possesses an improved structure integrity and higher specific surface area than the spherical shape, and exhibits extraordinary electrochemical performance. Therefore, the strategy of developing hierarchical morphologies with preferentially oriented active planes is promising for improving the performance of LLO materials.

5.1. Preferentially exposed crystallographic facets The surface structure is an essential aspect that controls the Li+ diffusion kinetics in the crystal lattice. It has been established that Li+ can merely diffuse along the Li layers in LLO materials with a αNaFeO2 structure. As shown in Fig. 17a [57], as the growth direction of an LLO nanoplate is vertical to [001], the surface shows (001) planes preferentially that are inactive electrochemically since no appropriate path is offered for Li+ diffusion [94]. On the contrary, the (010), (110) and (100) planes that are normal to the (001) plane are considered to be active for Li+ diffusion [55]. Based on theoretical calculations, Sun's group indicated that the high-energy (010) plane is unfavorable energetically during growth process, when compared with the (001) plane [57]. Therefore, (001) nanoplates are thermodynamically stable products in LLO materials in various production routes [95–98]. By adjusting the synthesis conditions, Sun's group reported a crystal habit-tuned nanoplate LLO with preferentially growth of (010) planes (Fig. 17b), when compared with the conventional thermodynamically stable material [57]. To integrate the merits of the hierarchical structure and the preferentially exposed (010) planes of the LLOs, Wu's group developed a strategy to synthesize a hierarchical Li1.2Mn0.6Ni0.2O2 cathode material with increased surface area of exposed {010} planes (Fig. 17c and d) [59]. This material exhibits excellent rate performance and good capacity retention. Following this research, they reported the synthesis of hierarchical Li1.2Mn0.54Ni0.13Co0.13O2 spheres using an ionic interfusion method. The hierarchical LLO spheres were assembled with primary nanoplates with increased amounts of electrochemically active planes, such as (110) planes (Fig. 17e and f) [99], which favor rapid Li+ intercalation/deintercalation, and the microscale spherical secondary particles with mesopores ensure structural stability and facilitate electrolyte transport. More recently, they demonstrated the synthesis of microscale fusiform Li1.2Ni0.2Mn0.6O2 cathodes assembled with nanoparticles, with preferentially oriented active (110) planes, using a hydrothermal method, in which urea and polyvinyl pyrrolidone were employed to guide the crystal growth by altering the

5.2. Electrically conductive coatings The low rate capability of LLOs is attributed to the insulative nature of the Li2MnO3 constituent and a high interfacial resistance along the LLO/electrolyte interfaces that is related to the SEI layer formed in high potential oxidation [39,41]. Various approaches have been developed to improve the rate performance of LLO materials. Amongst these approaches, conductive coatings based on carbonaceous material, such as carbon [101], graphene [102,103], carbon nanotubes [104] etc., and conductive polymers, such as polyaniline [105], polythiophene-based polymer [106] etc., can boost the rate capability by: (1) improving charge transfer kinetics across the LLO/electrolyte interfaces and (2) facilitating electron conduction among the active LLO particles and (3) providing additional paths for electron conduction between the LLO particles and the current collector. More recently, composite coatings including Li3PO4/C composite [107], and reduced graphene oxide (rGO)/AlPO4 composites [106] have been employed to enhance both the electronic and ionic conductivities of LLOs. Fig. 18a and b show the typical morphology of an rGO/AlPO4 coated LLO and the improved rate capability, respectively [108]. Carbon-based coatings usually involve pyrolytic process of a carbonaceous source on the LLO surface, which may produce reductive species like H2 and CO. The reductive species will reduce the TM cations and in turn change the crystal structure of the surface and subsurface regions of the LLOs. For instance, Wei's group employed an in-situ synchronous carbonization reduction process to develop a Layered@Spinel@Carbon heterostructure, which contains a layered core, a cubic spinel interlayer, and a nanoscaled carbon coating 594

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Fig. 19. Typical HRTEM and HAADF-STEM micrographs and EDX spectra taken from pristine (a, d and g), Al2O3 coated (b, e and h) and TiO2 coated (c, f and i) Li1.2Mn0.54Ni0.13Co0.13O2 materials prepared by ALD technique. (Reprinted with permission [120]. Copyright 2013 WILEY-VCH, Verlag GmbH & Co. KGaA.).

conductors possess relatively high Li conductivity with ion transport channels on the LLO surface and are inactive electrochemically in charge-discharge process. A ionic conductor Li3VO4-coated Li1.18Co0.15Ni0.15Mn0.52O2 cathode material has been developed via direct reaction with NH4VO3 at 350 °C [110]. The coated substrate material has a layered structure with reduced Li+/Ni2+ cation mixing, exhibiting a higher discharge capacity (or less IRC loss). The Li3VO4 coating also prohibits the growth of the SEI layer, thus enhancing the rate capability, thermal stability, and electrochemcial cyclability of LLOs simultaneously. Hu's group fabricated Li1.13Ni0.30Mn0.57O2 materials coated with fast Li+-ion conductor Li2SiO3 structure [113]. With the assistance of the distinct Li2SiO3 surface coating, the LLO material demonstrates superior rate performance and long-term cyclability, when compared with the pristine material. It was also shown that a Ni-doped Li3PO4 coating could significantly improve the rate capability of LLO materials by stabilizing the surface structure [111,112]. The improved rate capability can be associated with a defect Ni-doped

(Fig. 18c and d) [109]. The distinct heterostructure, integrating the merits of the high specific capacity of LLO structure, the high conductivity of the carbon coating, and fast Li+ diffusion kinetics of the spinel, shows high discharge capacity and excellent rate performance. It is worth noting, however, that deterioration in the reversible capacity and Coulombic efficiency would be prominent if too much layered structure was consumed to form the spinel structure. Ongoing effort is required to exploit a suitable coating technique that may be applied to produce carbonaceous coatings without a substantial destructive influence on the LLO materials. 5.3. Fast Li+-conducting coatings Another important strategy to address the limited rate capability by boosting the ionic conduction of Li cations is to apply ionic conductors as coating materials; these include Li3VO4 [110], Ni-doped Li3PO4 [111,112], Li2SiO3 [113], and LiPON [114]. Generally, these ionic 595

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Fig. 20. (a) The initial charge–discharge profiles of layered Li(Li0.2Mn0.54Ni0.13Co0.13)O2-V2O5 composite with different V2O5 contents. (Reprinted with permission [122]. Copyright 2008 Elsevier B.V.) (b)–(d) TEM morphology, schematic diagram, and cycling performance of MnO2 nanosheet coated Li1.2Mn0.567Ni0.167Co0.066O2. (Reprinted with permission [125]. Copyright 2014, The Royal Society of Chemistry.).

Li3PO4 surface structure, which promotes fast Li+-ion conduction at the LLO surface. Nanda's group shows the application of a nanoscale solid electrolyte of lithium phosphorus oxynitride (LiPON) with fast Li conductivity as a coating layer for the LLOs [114]. The LiPON-coated LLO cathode exhibited excellent rate capability and capacity retention, which is ascribed to the impact of LiPON nanolayer on facilitating the electrochemical activity and enhancing the interfacial stability along the LLO/electrolyte interface [114].

enhancement could be ascribed to the initial layered-to-spinel structure transformation, which is probably caused by the Li-percolating effect of the AlF3 coating. On the basis of detailed STEM and EELS analyses, Wang's group revealed that in addition to a layered to spinel transition, the AlF3 coating can greatly reduce formation of the SEI layer and protect the electrode from attack by HF in the electrolyte [118]. Manthiram's group investigated the effects of surface coatings with ZnO, Al2O3, ZrO2, SiO2, and CeO2 on LLOs [119]. When compared with the pristine material, less first cycle IRC loss and superior discharge capacity are observed in the coated LLO cathodes. The improvement is due to portion of the oxygen vacancies generated by the irreversible oxygen evolution in the first charge process being reserved in the LLO lattice. Surface coating with Al2O3 was discovered to be particularly efficient in maintaining more oxygen vacancies and restraining IRC loss. Atomic layer deposition (ALD) technique was also employed to prepare nanoscaled Al2O3 and TiO2 coatings on the LLO cathodes [120]. The Al2O3 coating prepared by the ALD technique exhibited a conformal morphology, whereas a rough morphology was observed in the TiO2 coating, as shown in Fig. 19 [120]. The Al2O3 coating was found to be stable upon extended cycles and showed enhanced cyclability, whereas the particulated TiO2 coating was more vulnerable to form a LixTiO2 interlayer. Therefore, the cyclic intercalation/deintercalation processes of Li+ cations leads to corrosion of the TiO2 coating, which accounts for performance degradation in LLO cathodes. Wang's group further revealed that in terms of surface chemical function, the Al2O3 coating refrains the side reactions along the cathode/electrolyte interfaces during battery cycling [121]. Meanwhile, the Al2O3 coating layer also eliminates the chemical reduction of Mn from the cathode surface, thereby preventing the dissolution of reduced Mn into the electrolyte [121]. In terms of structural stability, they found that the Al2O3 coating layer could

5.4. Electrochemically inactive protective coatings LiPF6 is the major lithium salt commercially available for LIBs and is rather sensitive to moisture to generate a side-product of hydrogen fluoride (HF). In an LLO system cycled to high potentials over 4.6 V, carbonate-based solvents are oxidized by the generated oxygen radicals. Water is one of the byproducts, leading to more HF generated in the LLO system when compared with other Li-ion systems. HF has generally been responsible for chemical dissolution of TM cations and chemical attack of LLO cathode surfaces [52]. When appropriate coating materials are applied to the LLO cathodes, they may act as an effective HF scavenger to form metal florides and neutralize the acidity near the electrode/electrolyte interfaces [115]. The newly formed metal fluorides are usually insoluble in carbonate-based electrolytes and practically serve as a protecting coating to resrtain the chemical attack on LLO surfaces. Therefore, it is reasonably expected that the incorporation of an HF scavenger can enhance the capacity retention of LLO cathodes to some extent. Coating materials for this purpose include ZnO, TiO2, Al2O3, ZrO2, SiO2, CeO2, and AlF3. AlF3-coated Li(Li0.19Ni0.16Co0.08Mn0.57)O2 cathodes with significant enhancement of the first cycle IRC, long-term cyclability, and thermal stability, have been developed [116,117]. The 596

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Fig. 21. (a)–(c) TEM and HRTEM images and enhanced cyclability of LiFePO4-coated LLOs. (Reprinted with permission [126]. Copyright 2013 WILEY-VCH, Verlag GmbH & Co. KGaA.) (d) and (e) HRTEM image and rate-dependent cycle performance of nanoscaled spinel coatd-LLO cathode. (Reprinted with permission [128]. Copyright 2014 American Chemical Society.).

will be detrimental for retaining the high discharge voltage of LLOs (Fig. 20a). Zhou's group employed delaminated MnO2 nanosheets to coat LLO materials to achieve comprehensive enhancement in electrochemical properties such as the initial coulombic efficiency, reversible discharge capacity, rate performance, and cyclability (Fig. 20b–d) [125]. These improvements are ascribed to the MnO2 nanosheet layer with enhanced intercalation activity and lower charge transfer resistance [125]. Note that the introduction of Li-insertion host materials can mitigate the IRC loss but does not resolve the root cause of irreversible oxygen evolution and subsequent TM migration in LLOs. Another class of electrochemically active materials that are applied to modify the LLO surface includes LiFePO4 (LFP) [126] and spineltype oxide materials [127–129]. Liu's group described a sol–gel process to incorporate nanoscaled LFP coating to the LLOs [126]. In addition to provide some reversible capacity, the thin LiFePO4 coating acts as a protective layer to protect the LLO substrate from chemical attack by electrolytes. Moreover, Liu's group claimed that a small amount of Fe and P was doped into the LLO structure that may stabilize the LLO lattice (Fig. 21a–c) [126]. Wu's group proposed a biomimetic strategy to synthesize an ultrathin spinel membraneencapsulated LLO cathode (Fig. 21d and e) [128]. The nanoscaled spinel film is responsible for the high reversible capacity, excellent rate performance, and superior cyclability of the LLO material [128]. Li's group reported a “mixed-oxalate” method to synthesize a layered/ spinel heterostructure through coating a LLO material with a LiNixMn2–xO4 spinel nanolayer [127]. This heterostructure inherits

mitigate the layer to spinel phase transformation during battery cycling [121]. The corrosion reaction induced by HF is a slow and steady procedure, so HF will gradually consume the protective coatings. Therefore, the effectiveness of thin protective coatings needs to be verified over a long-term cycling or through aging tests. Also, it is noted that perfect coverage using a protective coating is essential to reduce the unwanted site reactions along the cathode/electrolyte interfaces, but it is not necessary for the coating materials to serve as an effective HF scavenger. Consequently, coating technology needs to be optimized to achieve desirable long-term performance.

5.5. Multi-functional electrochemically active coatings The large IRC loss in the first cycle that is linked to the unique first charge process is another crucial obstacle for the commercial application of LLOs. In order to suppress the large IRC loss, the application of a Li-intercalation host material to capture the excess Li cations is an effective approach. Manthiram's group demonstrated that the first cycle IRC loss could be significantly suppressed through manipulating the content of Li-insertion hosts such as V2O5, LiV3O8, Li4Mn5O12, and VO2(B) [122–124]. For instance, the IRC loss reduces from 68 mA h g−1 for pure Li(Li0.2Mn0.54Ni0.13Co0.13)O2 to 0 mA h g−1 at around 89 wt% Li(Li0.2Mn0.54Ni0.13Co0.13)O2−11 wt% V2O5 since V2O5 serves as an intercalation host to hold the excess Li cations after the first charge process (Fig. 20a) [122]. Note that too much V2O5 content 597

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Fig. 22. (a) HAADF-STEM image, (b) ABF-STEM image, and (c) ABF contrast profile taken from Na+ doped LLO along the [010] zone axis. Rate-dependent performance (d) and cyclability (e) of pristine and Na+ doped LLO materials. (Reprinted with permission [139]. Copyright 2015 WILEY-VCH, Verlag GmbH & Co. KGaA.).

5.7. Summary of surface modification on the LLO materials

both advantages of the high reversible capacity of the layered structure and fast Li+ insertion/extraction kinetics of the spinel structure.

Surface modifications including chemical treatments, coatings, additives, or gradient dopants have been presented to enhance the reversible capacity, cyclability, and rate capability to some extent, but the voltage decay phenomenon remains upon long-term cycling, suggesting that TM migration and subsequent phase transformation also take place in the bulk material. Gradient surface doping strategies and the design of full concentration gradient LLOs, which integrate the advantages of surface protective layers and composition manipulation of bulk materials, are currently being intensely pursued to reduce voltage decay in LLOs. Nevertheless, continued efforts at elucidating the structural and chemical complexities along the modified electrode/ electrolyte interfaces at atomic scale would shed more light on the specific roles of surface modifications, and in turn, to provide some guidelines to optimize both synthesis procedures and electrochemical performances.

5.6. Gradient surface doping layers Chemical doping, including Mg [130], Al [131], Ti [132], Sn [133], Ru [93], Y [134], Zn [135], B [136], S [137], P [138] etc., has been considered to have positive effects on the cyclic durability and rate performance of LLOs. However, a homogeneous doping strategy is not sufficient to shield the oxide surface from unwanted chemical reactions with the electrolytes. Recently, Guo's group proposed a gradient surface Na+ doping strategy to realize an important pinning effect to stabilize the layered structure and to facilitate the Li+ diffusion kinetics in the LLO structure [139]. The Na+ gradient doping is achieved by calcinating the LLO materials in molten NaCl, which is thermodynamically driven by Na+ concentration diffusion. Consequently, the surface modified large particle LLO materials exhibit high Coulombic efficiency and remarkable cyclability without sacrificing high reversible capacity (Fig. 22) [139]. More recently, Wei's group developed a polyanion-doping scheme to originate structural transformation to the LLO surface, which generates a surface spinel layer and a gradient-doped layered bulk simultaneously (Fig. 23) [140]. The polyanions include BO45−, SiO44−, and PO43− etc. This gradient-doping method combines the benefits of bulk doping that stabilizes the layered lattice, and surface modification that protects the cathode surface from chemical attack by electrolytes [140]. A phosphate modified LLO material exhibits a high reversible capacity of about 300 mA h g−1, better electrochemical kinetics, and outstanding cyclability. Moreover, they found that in borate-doped LLOs treated at lower temperatures (~700 oC), layered NiO-type rock salt and spinel-type structures were present within a LLO primary particle, which is closely associated with the borate gradient. Optimal annealing temperature and borate content are considered to ensure the enhancement in cyclability and rate performance, indicative of the feasibility of low temperature synthesis of LLO materials [141].

6. Concluding remarks Lithium-rich layered oxides (LLOs) are a class of promising candidates for LIB cathodes because of their high discharge capacity of 250–300 mA h g−1 and high energy density of ~900 W h kg−1 at the cathode level, showing the potential to meet the target of 250 Wh kg−1 at the cell level for applications in electric vehicles and distributed energy storages. However, the challenges such as continuous voltage decay, poor rate performance, and structural instability remain, and substantial effort would be necessary to overcome these challenges. Over the past decade, with the assistance of aberration-corrected scanning transmission electron microscopy (STEM), coupled with electron energy loss spectroscopy (EELS) and energy dispersive X-ray (EDX) spectroscopy, and X-ray photoelectron spectroscopy (XPS), significant advancement has been achieved to appreciate the distinct surface structures in pristine materials and surface structural changes that accounts for the first cycle IRC and fast voltage/capacity fading. Several strategies have also been developed to alter the surface 598

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Fig. 23. Typical HAADF-STEM images, contrast profiles, and structural model along the same zone axis of (a) pristine material and (b) phosphate-modified material. (c) The first charge–discharge curves and rate-dependent cyclability of pristine and phosphate-modified LLOs. (Reprinted with permission [140]. Copyright 2016 WILEY-VCH, Verlag GmbH & Co. KGaA.).

tion parameters and chemistry is urgently needed. Electron and X-ray based tomography techniques can be used to develop a better understanding of LLO facet formation and the orientation relative to the core material, which will be useful to elucidate the corresponding effects on delithiation and cycling. Specifically, how reversible oxide-ion oxidation, the process of TM and Li migration, and generated O2 gas and oxide radicals are affected by surface structure and chemistry is of crucial importance. Ultimately, such findings can guide researchers in tailoring exposed facets on the surface, which can lessen the structural evolution and, thus, the cyclic stress and strain in LLOs upon cycling. The reversible and irreversible redox processes of lattice oxygen play a key role in electrochemical performance, including structural stability and electron and ion transport kinetics of LLO cathodes. Further effort on developing approaches to reduce irreversible oxygen evolution and effectively utilize the distinct activity of reversible oxygen redox in LLO cathodes is crucial in enhancing cycling stability and rate

structure and chemistry to tackle these critical aspects. For instance, novel approaches have been proposed to preparation crystal habittuned materials and hierarchical structures that can alleviate the IRC loss, voltage decay, and sluggish kinetics. Surface treatments such as coatings or gradient doping may improve the electrochemical properties of LLOs by stabilizing the bulk structure and shielding the surface from side reactions with electrolytes. The approaches mentioned in this review have alleviated the technical challenges of LLOs to some extent; however, substantial improvements in electrode kinetics and electrochemical cyclability through reasonable and economical approaches are still necessary for these materials to reach commercial implementation. Various studies have explored the surface structure, but no consensus has been reached, which is probably owing to the various manufacturing processes and compositions applied in different research. Therefore, a better understanding of the surface structure in relation to produc-

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Mater. 24 (2012) 3558–3566. [21] K.A. Jarvis, Z.Q. Deng, L.F. Allard, A. Manthiram, P.J. Ferreira, Chem. Mater. 23 (2011) 3614–3621. [22] B. Ammundsen, J. Paulsen, I. Davidson, R.-S. Liu, C.-H. Shen, J.-M. Chen, L.Y. Jang, J.-F. Lee, J. Electrochem. Soc. 149 (2002) A431–A436. [23] K.A. Jarvis, C.C. Wang, A. Manthiram, P.J. Ferreira, J. Mater. Chem. A 2 (2014) 1353–1362. [24] C.C. Wang, K.A. Jarvis, P.J. Ferreira, A. Manthiram, Chem. Mater. 25 (2013) 3267–3275. [25] E. McCalla, A.W. Rowe, R. Shunmugasundaram, J.R. Dahn, Chem. Mater. 25 (2013) 989–999. [26] E. McCalla, C.M. Lowartz, C.R. Brown, J.R. Dahn, Chem. Mater. 25 (2013) 912–918. [27] E. McCalla, A.W. Rowe, C.R. Brown, L.R.P. Hacquebard, J.R. Dahn, J. Electrochem. Soc. 160 (2013) A1134–A1138. [28] E. McCalla, J. Li, A.W. Rowe, J.R. Dahn, J. Electrochem. Soc. 161 (2014) A606–A613. [29] A.R. Armstrong, M. Holzapfel, P. Novak, C.S. Johnson, S.H. Kang, M.M. Thackeray, P.G. Bruce, J. Am. Chem. Soc. 128 (2006) 8694–8698. [30] C.S. Johnson, J.S. Kim, C. Lefief, N. Li, J.T. Vaughey, M.M. Thackeray, Electrochem. Comm. 6 (2004) 1085–1091. [31] J.M. Zheng, Z.R. Zhang, X.B. Wu, Z.X. Dong, Z. Zhu, Y. Yang, J. Electrochem. Soc. 155 (2008) A775–A782. [32] J. Hong, H.D. Lim, M. Lee, S.W. Kim, H. Kim, S.T. Oh, G.C. Chung, K. Kang, Chem. Mater. 24 (2012) 2692–2697. [33] N. Tran, L. Croguennec, M. Menetrier, F. Weill, P. Biensan, C. Jordy, C. Delmas, Chem. Mater. 20 (2008) 4815–4825. [34] J. Hong, D.H. Seo, S.W. Kim, H. Gwon, S.T. Oh, K. Kang, J. Mater. Chem. 20 (2010) 10179–10186. [35] B. Xu, C.R. Fell, M.F. Chi, Y.S. Meng, Energy Environ. Sci. 4 (2011) 2223–2233. [36] T. Nakajima, H. Groult, Fluorinated Materials for Energy Conversion, Elsevier, Oxford, UK, 2005. [37] E. Peled, J. Electrochem. Soc. 126 (1979) 2047–2051. [38] E. Peled, D. Golodnitsky, G. Ardel, J. Electrochem. Soc. 144 (1997) L208–L210. [39] J. Yan, X. Liu, B. Li, RSC Adv. 4 (2014) 63268–63284. [40] H. Yu, H. Zhou, J. Phys. Chem. Lett. 4 (2013) 1268–1280. [41] A. Manthiram, J.C. Knight, S.T. Myung, S.-M. Oh, Y.-K. Sun, Adv. Energy Mater. 6 (2016) 1501010. [42] P. Yan, J. Zheng, J. Xiao, C. Wang, J.G. Zhang, Front. Energy Res. 3 (2015). http://dx.doi.org/10.3389/fenrg.2015.00026. [43] D. Mohanty, J. Li, S.C. Nagpure, D.L.I. Wood, C. Daniel, MRS Energy Sustain. 2 (2015). http://dx.doi.org/10.1557/mre.2015.16. [44] S. Hy, H. Liu, M. Zhang, D. Qian, B.J. Hwang, Y.S. Meng, Energy Environ. Sci. 9 (2016) 1931–1954. [45] H. Moriwake, A. Kuwabara, C.A.J. Fisher, R. Huang, T. Hitosugi, Y.H. Ikuhara, H. Oki, Y. Ikuhara, Adv. Mater. 25 (2013) 618–622. [46] K. Kang, Y.S. Meng, J. Breger, C.P. Grey, G. Ceder, Science 311 (2006) 977–980. [47] T. Ohzuku, Y. Makimura, Chem. Lett. 30 (2001) 744–745. [48] J. Lee, A. Urban, X. Li, D. Su, G. Hautier, G. Ceder, Science 343 (2014) 519–522. [49] A. Urban, J. Lee, G. Ceder, Adv. Energy Mater. 4 (2014) 1400478. [50] Y. Lyu, N. Zhao, E. Hu, R. Xiao, X. Yu, L. Gu, X.Q. Yang, H. Li, Chem. Mater. 27 (2015) 5238–5252. [51] H. Dixit, W. Zhou, J.C. Idrobo, J. Nanda, V.R. Cooper, ACS Nano 8 (2014) 12710–12716. [52] M. Gu, I. Belharouak, A. Genc, Z.G. Wang, D.P. Wang, K. Amine, F. Gao, G.W. Zhou, S. Thevuthasan, D.R. Baer, J.G. Zhang, N.D. Browning, J. Liu, C.M. Wang, Nano Lett. 12 (2012) 5186–5191. [53] J. Zheng, M. Gu, A. Genc, J. Xiao, P. Xu, X. Chen, Z. Zhu, W. Zhao, L. Pullan, C. Wang, J.G. Zhang, Nano Lett. 14 (2014) 2628–2635. [54] P.F. Yan, J.M. Zheng, J.X. Zheng, Z.G. Wang, G.F. Teng, S. Kuppan, J. Xiao, G.Y. Chen, F. Pan, J.G. Zhang, C.M. Wang, Adv. Energy Mater. 6 (2016) 1502455. [55] H. Chen, C.P. Grey, Adv. Mater. 20 (2008) 2206–2210. [56] D.H. Park, S.T. Lim, S.J. Hwang, C.S. Yoon, Y.K. Sun, J.H. Choy, Adv. Mater. 17 (2005) 2834–2837. [57] G.Z. Wei, X. Lu, F.S. Ke, L. Huang, J.T. Li, Z.X. Wang, Z.Y. Zhou, S.G. Sun, Adv. Mater. 22 (2010) 4364–4367. [58] F. Fu, G.L. Xu, Q. Wang, Y.P. Deng, X. Li, J.T. Li, L. Huang, S.G. Sun, J. Mater. Chem. A 1 (2013) 3860–3864. [59] L. Chen, Y.F. Su, S. Chen, N. Li, L.Y. Bao, W.K. Li, Z. Wang, M. Wang, F. Wu, Adv. Mater. 26 (2014) 6756–6760. [60] H. Koga, L. Croguennec, M. Menetrier, K. Douhil, S. Belin, L. Bourgeois, E. Suard, F. Weill, C. Delmas, J. Electrochem. Soc. 160 (2013) A786–A792. [61] H. Koga, L. Croguennec, M. Ménétrier, P. Mannessiez, F. Weill, C. Delmas, J. Power Sources 236 (2013) 250–258. [62] M. Sathiya, K. Ramesha, G. Rousse, D. Foix, D. Gonbeau, A.S. Prakash, M.L. Doublet, K. Hemalatha, J.M. Tarascon, Chem. Mater. 25 (2013) 1121–1131. [63] M. Sathiya, G. Rousse, K. Ramesha, C.P. Laisa, H. Vezin, M.T. Sougrati, M.L. Doublet, D. Foix, D. Gonbeau, W. Walker, A.S. Prakash, M. Ben Hassine, L. Dupont, J.M. Tarascon, Nat. Mater. 12 (2013) 827–835. [64] C.S. Johnson, N. Li, C. Lefief, M.M. Thackeray, Electrochem. Commun. 9 (2007) 787–795. [65] A.D. Robertson, P.G. Bruce, Chem. Commun. (2002) 2790–2791. [66] N. Yabuuchi, K. Yoshii, S.T. Myung, I. Nakai, S. Komaba, J. Am. Chem. Soc. 133 (2011) 4404–4419. [67] C. Genevois, H. Koga, L. Croguennec, M. Ménétrier, C. Delmas, F. Weill, J. Phys. Chem. C 119 (2015) 75–83.

performance. For instance, Meng's group demonstrated a gas-solid interface reaction to obtain a controlled amount of oxygen vacancies in the surface region, which substantially suppress irreversible oxygen gas release and subsequent voltage/capacity fading [142]. In this respect, the development of operando structural characterization and spectroscopy techniques to probe lattice oxygen and TM cations will be important. In order to acquire structural and chemical information from the surface to the bulk, complementary high-resolution electron microscopy and spectroscopy techniques probing at multiple length scales will be necessary for material evaluation. Full cell studies may offer insight into the actual interfacial phenomena occurred to the electrode/electrolyte interfaces; therefore, full-cell considerations based on the LLO cathode materials would be crucial. Continued studies on various LLO compositions and new electrolytes will also be required to establish the material propertyelectrochemical performance correlation and to offer reasonable design guidelines. A new area of surface modification in LLO research is the development of new electrolyte additives, which could lead to more stable and safer electrodes by controlled decomposition of electrolytes. If electrolyte additives are electroactive and act as the adsorbent for oxygen radicals, undesirable side reactions on the electrode/electrolyte interfaces will be suppressed, and in turn, considerably improved electrochemical properties of LLOs are expected. Further investigation is still needed in this research field. Acknowledgements The authors would like to acknowledge the financial support from the Recruitment Program of Global Youth Experts, the National Natural Science Foundation of China (51304248), the Program for New Century Excellent Talents in University (NCET-11-0525), the Doctoral Fund of Ministry of Education of China (20130162110002), the Innovation Program of Central South University (2016CXS003) and the State Key Laboratory of Powder Metallurgy at Central South University. References [1] M.M. Thackeray, C. Wolverton, E.D. Isaacs, Energy Environ. Sci. 5 (2012) 7854–7863. [2] M.M. Thackeray, C.S. Johnson, J.T. Vaughey, N. Li, S.A. Hackney, J. Mater. Chem. 15 (2005) 2257–2267. [3] B.L. Ellis, K.T. Lee, L.F. Nazar, Chem. Mater. 22 (2010) 691–714. [4] T. Ohzuku, M. Nagayama, K. Tsuji, K. Ariyoshi, J. Mater. Chem. 21 (2011) 10179–10188. [5] M.M. Thackeray, S.-H. Kang, C.S. Johnson, J.T. Vaughey, R. Benedek, S.A. Hackney, J. Mater. Chem. 17 (2007) 3112–3125. [6] F. Zhou, X. Zhao, A. van Bommel, X. Xia, J.R. Dahn, J. Electrochem. Soc. 158 (2011) A187–A191. [7] M.H. Rossouw, D.C. Liles, M.M. Thackeray, J. Solid State Chem. 104 (1993) 464–466. [8] Y. Wu, A. Manthiram, J. Power Sources 183 (2008) 749–754. [9] M.M. Thackeray, S.H. Kang, C.S. Johnson, J.T. Vaughey, S.A. Hackney, Electrochem. Commun. 8 (2006) 1531–1538. [10] H. Yu, H. Kim, Y. Wang, P. He, D. Asakura, Y. Nakamura, H. Zhou, Phys. Chem. Chem. Phys. 14 (2012) 6584–6595. [11] J.G. Wen, J. Bareno, C.H. Lei, S.H. Kang, M. Balasubramanian, I. Petrov, D.P. Abraham, Solid State Ion. 182 (2011) 98–107. [12] H.J. Yu, Y.R. Wang, D. Asakura, E. Hosono, H.S. Zhou, RSC Adv. 2 (2012) 8797–8807. [13] H.J. Yu, R. Ishikawa, Y.G. So, N. Shibata, T. Kudo, H. Zhou, Y. Ikuhara, Angew. Chem. Int. Ed. 52 (2013) 5969–5973. [14] H. Koga, L. Croguennec, P. Mannessiez, M. Ménétrier†, F. Weill, L. Bourgeois, M. Duttine, E. Suard, C. Delmas, J. Phys. Chem. C 116 (2012) 13497–13506. [15] K. Jarvis, Z. Deng, A. Manthiram, P.J. Ferreira, L.F. Allard, Microsc. Microanal. 18 (2012) 1414–1415. [16] K. Jarvis, Z. Deng, A. Manthiram, P.J. Ferreira, Microsc. Microanal. 18 (2012) 1484–1485. [17] A.K. Shukla, Q.M. Ramasse, C. Ophus, H. Duncan, F. Hage, G.Y. Chen, Nat. Comm. 6 (2015) 8711–8719. [18] Z. Lu, Z. Chen, J.R. Dahn, Chem. Mater. 15 (2003) 3214–3220. [19] J. Hong, H. Gwon, S.-K. Jung, K. Ku, K. Kang, J. Electrochem. Soc. 162 (2015) A2447–A2467. [20] A. Boulineau, L. Simonin, J.F. Colin, E. Canevet, L. Daniel, S. Patoux, Chem.

600

Nano Energy 30 (2016) 580–602

W. Wei et al.

[114] S.K. Martha, J. Nanda, Y. Kim, R.R. Unocic, S. Pannala, N.J. Dudney, J. Mater. Chem. A 1 (2013) 5587–5595. [115] Z.H. Chen, J.R. Dahn, Electrochem. Solid-State Lett. 6 (2003) A221–A224. [116] Y.K. Sun, M.J. Lee, C.S. Yoon, J. Hassoun, K. Amine, B. Scrosati, Adv. Mater. 24 (2012) 1192–1196. [117] S. Sun, Y. Yin, N. Wan, Q. Wu, X. Zhang, D. Pan, Y. Bai, X. Lu, ChemSusChem 8 (2015) 2544–2550. [118] J. Zheng, M. Gu, J. Xiao, B.J. Polzin, P. Yan, X. Chen, C.M. Wang, J.G. Zhang, Chem. Mater. 26 (2014) 6320–6327. [119] Y. Wu, A. Manthiram, Solid State Ion. 180 (2009) 50–56. [120] X.F. Zhang, I. Belharouak, L. Li, Y. Lei, J.W. Elam, A.M. Nie, X.Q. Chen, R.S. Yassar, R.L. Axelbaum, Adv. Energy Mater. 3 (2013) 1299–1307. [121] P.F. Yan, J.M. Zheng, X.F. Zhang, R. Xu, K. Amine, J. Xiao, J.G. Zhang, C.M. Wang, Chem. Mater. 28 (2016) 857–863. [122] J. Gao, J. Kim, A. Manthiram, Electrochem. Commun. 11 (2009) 84–86. [123] J. Gao, A. Manthiram, J. Power Sources 191 (2009) 644–647. [124] E.S. Lee, A. Manthiram, J. Electrochem. Soc. 158 (2011) A47–A50. [125] S.H. Guo, H.J. Yu, P. Liu, X.Z. Liu, D. Li, M.W. Chen, M. Ishida, H.S. Zhou, J. Mater. Chem. A 2 (2014) 4422–4428. [126] F.H. Zheng, C.H. Yang, X.H. Xiong, J.W. Xiong, R.Z. Hu, Y. Chen, M.L. Liu, Angew. Chem. Int. Ed. 54 (2015) 13058–13062. [127] D. Luo, G. Li, C. Fu, J. Zheng, J. Fan, Q. Li, L. Li, Adv. Energy Mater. 4 (2014) 1400062. [128] F. Wu, N. Li, Y. Su, H. Shou, L. Bao, W. Yang, L. Zhang, R. An, S. Chen, Adv. Mater. 25 (2013) 3722–3726. [129] F. Wu, N. Li, Y.F. Su, L. Zhang, L. Bao, J. Wang, L. Chen, Y. Zheng, L. Dai, J. Peng, S. Chen, Nano Lett. 14 (2014) 3550–3555. [130] G.H. Kim, S.T. Myung, H.S. Kim, Y.K. Sun, Electrochim. Acta 51 (2006) 2447–2453. [131] S.H. Park, Y.K. Sun, J. Power Sources 119 (2003) 161–165. [132] V. Subramanian, G.T.K. Fey, Solid State Ion. 148 (2002) 351–358. [133] Y. Wang, Z. Yang, Y. Qian, L. Gu, H. Zhou, Adv. Mater. 27 (2015) 3915–3920. [134] S.F. Kang, H.F. Qin, Y. Fang, X. Li, Y.G. Wang, Electrochim. Acta 144 (2014) 22–30. [135] G.T.K. Fey, J.G. Chen, V. Subramanian, T. Osaka, J. Power Sources 112 (2002) 384–394. [136] B. Li, H. Yan, J. Ma, P. Yu, D. Xia, W. Huang, W. Chu, Z. Wu, Adv. Funct. Mater. 24 (2014) 5112–5118. [137] H.Z. Zhang, F. Li, G.L. Pan, G.R. Li, X.P. Gao, J. Electrochem. Soc. 162 (2015) A1899–A1904. [138] H.Z. Zhang, Q.Q. Qiao, G.R. Li, X.P. Gao, J. Mater. Chem. A 2 (2014) 7454–7460. [139] R.P. Qing, J.L. Shi, D.D. Xiao, X.D. Zhang, Y.X. Yin, Y.B. Zhai, L. Gu, Y.G. Guo, Adv. Energy Mater. 5 (2015) 1501914. [140] Y. Zhao, J.T. Liu, S.B. Wang, R. Ji, Q.B. Xia, Z.P. Ding, W.F. Wei, Y. Liu, P. Wang, D.G. Ivey, Adv. Funct. Mater. 26 (2016) 4760–4767. [141] J.T. Liu, S.B. Wang, Z.P. Ding, R.Q. Zhou, J.F. Zhang, L.B. Chen, W.F. Wei, J.T. Wang, ACS Appl. Mater. Interface 8 (2016) 18008–18017. [142] B. Qiu, M. Zhang, L. Wu, J. Wang, Y. Xia, D. Qian, H. Liu, S. Hy, Y. Chen, K. An, Y. Zhu, Z. Liu, Y.S. Meng, Nat. Commun. 7 (2016). http://dx.doi.org/10.1038/ ncomms12108.

[68] H. Koga, L. Croguennec, M. Ménétrier, P. Mannessiez, F. Weill, C. Delmas, S. Belin, J. Phys. Chem. C 118 (2014) 5700–5708. [69] S.A. Freunberger, Y.H. Chen, Z.Q. Peng, J.M. Griffin, L.J. Hardwick, F. Barde, P. Novak, P.G. Bruce, J. Am. Chem. Soc. 133 (2011) 8040–8047. [70] S.A. Freunberger, Y.H. Chen, N.E. Drewett, L.J. Hardwick, F. Barde, P.G. Bruce, Angew. Chem. Int. Ed. 50 (2011) 8609–8613. [71] Y. Chen, S.A. Freunberger, Z. Peng, F. Barde, P.G. Bruce, J. Am. Chem. Soc. 134 (2012) 7952–7957. [72] C.S. Johnson, N.C. Li, C. Lefief, J.T. Vaughey, M.M. Thackeray, Chem. Mater. 20 (2008) 6095–6106. [73] A. Ito, Y. Sato, T. Sanada, M. Hatano, H. Horie, Y. Ohsawa, J. Power Sources 196 (2011) 6828–6834. [74] S.H. Yu, T. Yoon, J. Mun, S. Park, Y.S. Kang, J.H. Park, S.M. Oh, Y.E. Sung, J. Mater. Chem. A 1 (2013) 2833–2839. [75] K.J. Carroll, D. Qian, C. Fell, S. Calvin, G.M. Veith, M.F. Chi, L. Baggetto, Y.S. Meng, Phys. Chem. Chem. Phys. 15 (2013) 11128–11138. [76] M. Oishi, C. Yogi, I. Watanabe, T. Ohta, Y. Orikasa, Y. Uchimoto, Z. Ogumi, J. Power Sources 276 (2015) 89–94. [77] S. Fukuzumi, K. Ohkubo, Chem. Eur. J. 6 (2000) 4532–4535. [78] M. Jiang, B. Key, Y.S. Meng, C.P. Grey, Chem. Mater. 21 (2009) 2733–2745. [79] Z.Q. Deng, A. Manthiram, J. Phys. Chem. C 115 (2011) 7097–7103. [80] J.R. Croy, D. Kim, M. Balasubramanian, K. Gallagher, S.H. Kang, M.M. Thackeray, J. Electrochem. Soc. 159 (2012) A781–A790. [81] M. Bettge, Y. Li, K. Gallagher, Y. Zhu, Q.L. Wu, W.Q. Lu, I. Bloom, D.P. Abraham, J. Electrochem. Soc. 160 (2013) A2046–A2055. [82] D. Mohanty, A.S. Sefat, J.L. Li, R.A. Meisner, A.J. Rondinone, E.A. Payzant, D.P. Abraham, D.L. Wood, C. Daniel, Phys. Chem. Chem. Phys. 15 (2013) 19496–19509. [83] D. Mohanty, J.L. Li, D.P. Abraham, A. Huq, E.A. Payzant, D.L. Wood, C. Daniel, Chem. Mater. 26 (2014) 6272–6280. [84] A. Boulineau, L. Simonin, J.F. Colin, C. Bourbon, S. Patoux, Nano Lett. 13 (2013) 3857–3863. [85] Z. Lu, J.R. Dahn, J. Electrochem. Soc. 149 (2002) A815–A822. [86] Y. Wu, A. Vadivel Murugan, A. Manthiram, J. Electrochem. Soc. 155 (2008) A635–A641. [87] P. Yan, A. Nie, J. Zheng, Y. Zhou, D. Lu, X. Zhang, R. Xu, I. Belharouak, X. Zu, J. Xiao, K. Amine, J. Liu, F. Gao, R. Shahbazian-Yassar, J.G. Zhang, C.M. Wang, Nano Lett. 15 (2015) 514–522. [88] P. Oh, M. Ko, S. Myeong, Y. Kim, J. Cho, Adv. Energy Mater. 4 (2014). http:// dx.doi.org/10.1002/aenm.201400631. [89] J. Zheng, P. Xu, M. Gu, J. Xiao, N.D. Browning, P. Yan, C.M. Wang, J.G. Zhang, Chem. Mater. 27 (2015) 1381–1390. [90] S.K. Jung, H. Gwon, J. Hong, K.Y. Park, D.H. Seo, H. Kim, J. Hyun, W. Yang, K. Kang, Adv. Energy Mater. 4 (2014). http://dx.doi.org/10.1002/ aenm.201300787. [91] J. Reed, G. Ceder, A. Van der Ven, Electrochem. Solid-State Lett. 4 (2001) A78–A81. [92] K. Kang, G. Ceder, Phys. Rev. B 74 (2006) 094105. [93] M. Sathiya, A.M. Abakumov, D. Foix, G. Rousse, K. Ramesha, M. Saubanere, M.L. Doublet, H. Vezin, C.P. Laisa, A.S. Prakash, D. Gonbeau, G. VanTendeloo, J.M. Tarascon, Nat. Mater. 14 (2015) 230–238. [94] J.B. Bates, N.J. Dudney, B.J. Neudecker, F.X. Hart, H.P. Jun, S.A. Hackney, J. Electrochem. Soc. 147 (2000) 59–70. [95] X.L. Liu, J.J. Wu, X.L. Huang, Z.W. Liu, Y. Zhang, M. Wang, R.C. Che, J. Mater. Chem. A 2 (2014) 15200–15208. [96] M.G. Kim, M. Jo, Y.-S. Hong, J. Cho, Chem. Commun. (2009) 218–220. [97] J. Cho, Y. Kim, M.G. Kim, J. Phys. Chem. C 111 (2007) 3192–3196. [98] X. Qian, X. Cheng, Z.Y. Wang, X.J. Huang, R. Guo, D.L. Mao, C.K. Chang, W.J. Song, Nanotechnology 20 (2009) 115608. [99] L.J. Zhang, N. Li, B.R. Wu, H.L. Xu, L. Wang, X.Q. Yang, F. Wu, Nano Lett. 15 (2015) 656–661. [100] Y. Li, Y. Bai, C. Wu, J. Qian, G.H. Chen, L. Liu, H. Wang, X.Z. Zhou, F. Wu, J. Mater. Chem. A 4 (2016) 5942–5951. [101] J. Liu, Q. Wang, B. Reeja-Jayan, A. Manthiram, Electrochem. Comm. 12 (2010) 750–753. [102] B. Song, M.O. Lai, Z. Liu, H. Liu, L. Lu, J. Mater. Chem. A 1 (2013) 9954–9965. [103] K.C. Jiang, X.L. Wu, Y.X. Yin, J.S. Lee, J. Kim, Y.G. Guo, ACS Appl. Mater. Interface 4 (2012) 4858–4863. [104] J. Mun, J.H. Park, W. Choi, A. Benayad, J.H. Park, J.M. Lee, S.G. Doo, S.M. Oh, J. Mater. Chem. A 2 (2014) 19670–19677. [105] Q. Xue, J. Li, G. Xu, H. Zhou, X. Wang, F. Kang, J. Mater. Chem. A 2 (2014) 18613–18623. [106] J. Lee, W. Choi, J. Electrochem. Soc. 162 (2015) A743–A748. [107] H. Liu, C. Chen, C. Du, X. He, G. Yin, B. Song, P. Zuo, X. Cheng, Y. Ma, Y. Gao, J. Mater. Chem. A 3 (2015) 2634–2641. [108] I.T. Kim, J.C. Knight, H. Celio, A. Manthiram, J. Mater. Chem. A 2 (2014) 8696–8704. [109] Q.B. Xia, X.F. Zhao, M.Q. Xu, Z.P. Ding, J.T. Liu, L.B. Chen, D.G. Ivey, W.F. Wei, J. Mater. Chem. A 3 (2015) 3995–4003. [110] Q. Fu, F. Du, X.F. Bian, Y.H. Wang, X. Yan, Y.Q. Zhang, K. Zhu, G. Chen, C.Z. Wang, Y.J. Wei, J. Mater. Chem. A 2 (2014) 7555–7562. [111] D. Shin, C. Wolverton, J.R. Croy, M. Balasubramanian, S.H. Kang, C.M.L. River, M.M. Thackeray, J. Electrochem. Soc. 159 (2011) A121–A127. [112] S. Kang, M.M. Thackeray, Electrochem. Commun. 11 (2009) 748–751. [113] E.Y. Zhao, X.F. Liu, H. Zhao, X.L. Xiao, Z.B. Hu, Chem. Commun. 51 (2015) 9093–9096.

Weifeng Wei is a professor in the State Key Laboratory of Powder Metallurgy at the Central South University (China). After receiving his Ph.D. in Materials Engineering from the University of Alberta (2009), He joined the Sadoway Group as a postdoctoral research associate at MIT (2009–2011). His research concerns materials development in the fields of energy storage, including materials and device development for electrochemical supercapacitors and rechargeable batteries, and electrochemical recycling of spent materials in non-aquesous electrolytes.

Libao Chen is a professor in the State Key Laboratory of Power Metallurgy at the Central South University (China). He received his PhD in Materials Physics and Chemistry from the Shanghai Institute of Microsystem and Information Technology, Chinese Academy of Sciences in 2007. His research focuses on high performance electrode materials and electrochemical energy storage systems, including Li-ion batteries, Na-ion batteries and supercapacitors.

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W. Wei et al. Prof. Anqiang Pan received his doctoral degree (2011) in Materials Physics and Chemistry from Central South University. He joined Prof. Guozhong Cao's group at University of Washington as a visiting student in 2008. Then, he worked at PNNL as a visiting scholar in Dr. JiGuang Zhang and Jun Liu's group (2009–2011). He joined Prof. Xiongwen (David) Lou's group at Nanyang Technological University as a research Fellow (2011– 2012). Thereafter, he worked at Central South University as a Sheng-Hua Professor in 2013. His current interests are on lithium/sodium ion batteries, and supercapacitors.

Douglas Ivey is a professor in the Department of Chemical and Materials Engineering at the University of Alberta (Canada). He received his Ph.D. in Engineering Materials from the University of Windsor (Canada) in 1985. His research focuses on applying high resolution microstructural characterization techniques to understanding the relationships between materials structure, properties and processing. Recent work has focused on developing electrochemical techniques to deposit thin films and thicker coatings for a range of electrochemical energy storage applications, including solid oxide fuel cells, supercapacitors and rechargeable batteries.

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