Layered ternary metal oxides: Performance degradation mechanisms as cathodes, and design strategies for high-performance batteries

Journal Pre-proofs Layered ternary metal oxides: performance degradation mechanisms as cathodes, and design strategies for high-performance batteries Lehao Liu, Meicheng Li, Lihua Chu, Bing Jiang, Lin Ruoxu, Zhu Xiaopei, Guozhong Cao PII: DOI: Reference:

S0079-6425(20)30019-0 https://doi.org/10.1016/j.pmatsci.2020.100655 JPMS 100655

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Progress in Materials Science

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Please cite this article as: Liu, L., Li, M., Chu, L., Jiang, B., Ruoxu, L., Xiaopei, Z., Cao, G., Layered ternary metal oxides: performance degradation mechanisms as cathodes, and design strategies for high-performance batteries, Progress in Materials Science (2020), doi: https://doi.org/10.1016/j.pmatsci.2020.100655

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Layered ternary metal oxides: performance degradation mechanisms as cathodes, and design strategies for high-performance batteries Lehao Liu1,2, Meicheng Li1,*, Lihua Chu1, Bing Jiang1, Lin Ruoxu2, Zhu Xiaopei2, and Guozhong Cao3,* 1State

Key Laboratory of Alternate Electrical Power System with Renewable Energy Sources, North China Electric Power University, Beijing 102206, China; 2Citic Guoan MGL Power Source Technology Co. Ltd., Beijing 102200, China; 3Department of Materials Science and Engineering, University of Washington, Seattle 98195, USA *To

whom correspondence should be addressed: [email protected] (M. Li) and [email protected] (G. Cao).

Abstract: Layered Li[NixCoyMz]O2 (M=Mn or Al, so-called NCM/NCA) ternary cathode materials have attracted a lot of intensive research efforts for high-performance lithium-ion batteries, because of their combined advantages with respect to energy density, production cost and environmental friendliness. However, those ternary metal oxides (especially Ni-rich) suffer from a few electrochemical cycling problems, such as strong capacity fading, severe voltage decay and safety issues. These problems are attributable mainly to the instability/irreversibility of the chemical composition, crystal structure and particle morphology, and the consequent undesirable physical/chemical processes during the synthesis and lithiation/delithiation processes. To circumvent these obstacles, a variety of strategies based on materials, electrode and electrolyte designs are investigated to effectively stabilize the NCM/NCA cathodes and to improve the electrochemical and thermal performance. This review scrutinizes the performance degradation mechanisms of the NCM/NCA materials and summarizes the recent advances in the materials, electrode and electrolyte levels by focusing on the relationships between the composition, structure, morphology, and properties. This paper intends to provide an easy entry for a comprehensive, systematic and deep understanding of the fundamentals, and offer a critical analysis and summary what have been done in the field and what are the challenges or hurdles to overcome. Key words: Ternary cathode material; Modification strategy; Structure stability; Electrochemical performance; Thermal stability; Massive application

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Graphic abstract:

Research highlights: (1). NCM/NCA (especially Ni-rich) materials have great application prospects in Li-ion batteries for electric vehicles due to their high energy density and low cost; (2). Instability of the composition/structure/morphology during cycling is the key reason for the poor electrochemical/thermal performance of the NCM/NCA materials; (3). Improvements in the cycling stability and safety need modification at NCM/NCA materials, electrode, and electrolyte levels; (4). The development trends about the NCM/NCA materials modification focus on hybrid coating, concentration-gradient design, single-crystal particle control, functional electrolyte additive, etc.

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Contents 1. Introduction ...................................................................................................................5 2. Composition/structure/morphology/properties of NCM/NCA materials......................8 2.1. Crystal structure and Li+ ion diffusion ...............................................................9 2.2. Roles and effects (structure/chemistry/property) of metal ions .......................10 2.3. Effects of the preparation process ....................................................................14 2.4. Other changes (side reactions and morphology) during cycling process ........15 2.5. Performance degradation mechanisms .............................................................18 2.6. Remarks............................................................................................................20 3. NCM/NCA material design.........................................................................................21 3.1. Composition design–body doping with other elements ...................................21 3.1.1. Cation substitution.................................................................................22 3.1.2. Anion substitution .................................................................................33 3.1.3. Co-substitution ......................................................................................36 3.1.4. Remarks.................................................................................................37 3.2. Structural design–surface coating and compositing with other materials ........37 3.2.1. Oxides....................................................................................................39 3.2.2. Fluorides ................................................................................................44 3.2.3. Phosphates .............................................................................................45 3.2.4. Inert Li host compounds (ionic conductors)..........................................47 3.2.5. Electronic conductors ............................................................................50 3.2.6. Dual/hybrid conductors .........................................................................52 3.2.7. Cathode materials (including concentration gradient) ..........................54 3.2.8. Compositing with multifunctional materials .........................................64 3.2.9. Remarks.................................................................................................65 3.3. Morphology design...........................................................................................67 3.3.1. Special morphologies (size/pore/shape/configuration control) .............67 3.3.2. Single-crystal cathode particle ..............................................................70 3.3.3. Remarks.................................................................................................73 3.4. Co-modification................................................................................................74 4. Electrode design ..........................................................................................................76

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4.1. Composition adjustion......................................................................................77 4.1.1. Conductive additives .............................................................................77 4.1.2. Binders...................................................................................................80 4.1.3. Blending with other cathode materials ..................................................82 4.2. Fabrication process regulation..........................................................................85 4.3. Remarks............................................................................................................86 5. Electrolyte design ........................................................................................................87 5.1. Li salts ..............................................................................................................88 5.2. Solvents ............................................................................................................89 5.3. Additives...........................................................................................................92 5.3.1. SEI film formation agents .....................................................................93 5.3.2. Acid scavengers.....................................................................................97 5.3.3. Flame retardants ....................................................................................98 5.3.4. Overcharge inhibitors ..........................................................................100 5.3.5. Multifunctional additives.....................................................................101 5.4. Remarks..........................................................................................................103 6. Summary and outlook................................................................................................104 Acknowledgments. ........................................................................................................108 References .....................................................................................................................117

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1. Introduction Under the great pressure of over depletion of fossil fuels and environmental concerns on pollution and global warming, researchers are developing efficient and green energy conversion and storage technologies [1]. Electric energy storage systems have attracted increasing interest since they allow the wide utilization of the renewable natural energy resources such as solar, wind, tide and geotherm [2, 3]. Thanks to their high energy density, long life cycle, light weight, compact design and environmental friendliness, lithium-ion batteries (LIBs) have been widely applied as power sources in many fields from consumer electronics (e.g., mobile phones and laptops) to automotive industry and stationary storage system during the past decade [4-6]. In spite of such success, higher energy/power densities and charge rate, longer calendar life, lower cost and higher safety are still highly desired [7-14]. Amongst the necessary components in LIBs, cathode materials play an important role in the capacity, operation voltage, cycling stability, and thermal stability [15-20]. Commercial cathode materials, such as layered LiCoO2, spinel LiMn2O4 and olivine LiFePO4 possess obvious merits and disadvantages (Table 1) [21-24]. LiNiO2 has an αNaFeO2 structure with the oxygen ions forming a cubic close-packed sublattice, which sparked a great interest by Dyer et al. in 1954 [25], because of its high theoretical capacity and low cost. But its commercialization was hindered by the difficulty in preparing high-quality samples due to the tendency to form Li-deficient compounds of Li1-zNi1+zO2 and the poor cycling performance and thermal stability [26, 27]. Firstprinciples calculations of the energy barriers show that O and Li vacancies in the crystal especially with high temperature processing would result in the Ni migration to the Li layer and the structure instability [28]. LiCoO2 was considered as the best cathode material in terms of capacity (usually ~150 mAh g–1, up to 180 mAh g–1 at 4.45 V vs. Li/Li+ with proper doping and surface modification), cycle stability, rate capability, tap density (2.8–3.0 g cm–3), ease of synthesis and handling for electrode fabrication. However, it is suitable mostly for small devices, owing to its high price and toxicity in respect to cobalt, and structural instability caused by the transition from hexagonal to monoclinic phase, Co dissolution and oxygen release at a highly oxidized state [29]. LiMn2O4 exhibits high safety and low cost, but the cycling performance (~120 mAh g–1) is poor due to the Mn dissolution in electrolyte and the Jahn-Teller phase change [30]. LiFePO4 displays high cycling stability and safety, but its specific capacity (~170 mAh g–1), operating voltage (~3.4 V vs. Li/Li+, thereafter) and tap density (1.0–1.7 g cm–3) are relatively low and its rate performance is poor because of the low ionic and electrical conductivities [31, 32]. Layered ternary cathode materials (LiNixCoyMzO2, where M=Mn or Al, so-called NCM/NCA) possess high storage capacity of ≥200 mAh g–1, high potential regions of ≥4.3 V, low cost, and environmental benignity [33-36]. The NCM/NCA ternary metal

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oxides with improved structural stability are also actually known as the doped LiNiO2 with Co and Mn/Al. First-principles density functional theory (DFT) calculations indicate that the Co and Mn ions prefer to exist in the same layer while the Li ions are preferentially located in the Li layer closest to the Co and Mn layers [37]. The transition metal (Tm) distribution can ease the stress from Tm-O bond length difference by the different effective charges of the transition metals. The crystals with transition metals and Li ions can also remain stable even under highly-delithiated states, leading to the enhanced stability of LiNiO2 upon the introduction of transition metals (Details of the roles of the transition metals will be discussed in Section 2). Liu et al. first reported the layered LiNi1-x-yCoxMnyO2 (0
LiCoO2

LiNiO2

LiNi0.8Co0.2O2

LiMn2O4

LiFePO4

NCM/NCA

Research time (year)

1980

1985

1992

1983

1997

1999

Theoretic/experimental

274/

275/

275/

148/

170/

~280/

capacity (mAh g–1)

140–155

≥150

≥170

100–120

90–165

155–220

Voltage plateau

3.8

3.8

3.8

4.1

3.4

3.7

Cycling stability (time)

≥500

–

–

≥500

≥2000

≥800

Tap density (g cm )

2.8–3.0

2.4–2.6

–

2.2–2.4

1.0–1.7

2.0–3.0

Cost

High

Medium

Medium

Low

Low

Medium

Safety

Low

Low

Medium

Medium

High

Medium

Toxicity

High

Medium

Medium

Low

Low

Medium

Synthesis

Easy

Difficult

Difficult

Easy

Medium

LIB application

Portable

Portable

EV

EV

EV,

devices

devices

(V vs. Li/Li ) +

–3

Medium massive

energy storage

Portable devices, EV

In spite of the great research progress made on the layered NCM/NCA materials, there remain some challenges, such as high first cycle irreversibility, poor rate capacity, strong capacity fading, significant decrease in discharge voltage plateau on the condition of successive cycling and safety issues related to the thermal runaway

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reactions. There are some excellent papers published on the recent advancement in the cathode materials especially for the Ni-rich and Li-rich materials [16, 40-51], and they always attributed the existed cycling problems to cation mixing and surface chemistry and summarized the partial progress on the modification methods of doping and coating in the materials level. In this review article, we take a different approach to disclose the performance degradation mechanisms of the NCM/NCA materials, emphasizing the relationships between the chemical composition, structure, morphology and properties of the layered NCM/NCA materials. Based on these analyses in the materials synthesis/storage, electrode preparation, cell fabrication and battery cycling processes, we ascribe the poor electrochemical and thermal performance of the NCM/NCA materials mainly to the changes (or instability) of the chemical composition (e.g., transition metal dissolution, Li depletion and oxygen evolution), structure (e.g., cation mixing and solid state interphase formation) and morphology (e.g., grain micro-crack and particle rupture), and the consequent physical/chemical phenomena (e.g., side reactions and heat generation) especially during the synthesis and lithiation/delithiation processes at high charge-discharge rates or temperatures (Fig. 1). Correspondingly, we scrutinize the up-to-date advance in the engineering strategies from the modification of composition (e.g., body doping), structure (e.g., surface coating and compositing) and morphology (e.g., size/shape/porosity/configuration and single-crystal particle) of the NCM/NCA ternary materials to the electrode design (e.g., conductive additives, binders, physical blending with other cathode materials, and adjustion of the whole cathode composition and preparation method), and then to the electrolyte design (e.g., novel Li salts, solvents, and functional additives). A scientific step-to-step story why the earlier researchers try to use these design methods and modification approaches linked with proper reasons is given in the article. Experimental studies and first-principles calculations are also combined to unravel the composition, structure and physical/electrochemical properties. This article intends to provide an easy entry for a comprehensive, systemic and profound understanding of all the above-mentioned NCM/NCA materials, in sharp contrast to the previously reported review/perspective/viewpoint/progress papers that could not articulately manifest the performance degradation mechanisms of the NCM/NCA cathodes and the essential design principles. We also outline the modification principles, preparation methods, composition/structure/morphology characteristics and physical/electrochemical properties of a few typical NCM/NCA materials, cathodes and electrolytes (Table 2). Furthermore, we give a few suggestions related to the development trends to better mature their battery applications. These principles and methods can be applied to other electrodes for improving their stability and achieving high-performance LIBs.

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Fig. 1 Engineering strategies of layered NCM/NCA materials for massive LIB applications.

2. Composition/structure/morphology/properties of NCM/NCA materials In Section 2, the chemical composition, crystal structure and morphology of the layered NCM/NCA cathode materials are discussed in detail, and also the effects of the preparation conditions and electrochemical charge-discharge (C-D) cycling on the cathode materials are summarized for better understanding the existed issues and challenges of the NCM/NCA materials, performance degradation mechanisms, and corresponding design principles and modification methods at the materials, electrode and electrolyte levels (in Section 3, 4 and 5) for high-performance LIBs.

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2.1. Crystal structure and Li+ ion diffusion Similar to the NCA materials, the NCM materials have an α-NaFeO2 layered _

structure with a R3m space group (Fig. 2A), which is a repeating O3 structure of oxygen-Li-oxygen-Tm-oxygen-Li-oxygen-Tm-oxygen along the rhombohedral [001] direction (Tm=Ni, Co and Mn). The Ni2+/3+, Co3+, Mn4+, Al3+ and Li+ ions are located in octahedral sites in a face-centered cubic oxygen stacked structure. For example, the Li, Tm and O atoms occupy the 3a, 3b and 6c sites in NCM333 materials, respectively, while the atoms of Ni, Co, Mn and Al, Li and O occupy the 3a, 3b and 6c sites in Nirich materials, respectively [45, 52]. A few factors such as the chemical composition, synthesis/storage conditions and electrochemical cycling process will also affect the ion occupy sites and the crystal structure (Discussed in section 2.2, 2.3 and 2.4).

Fig. 2 (A) Illustration of the ordered and disordered phases in layered Li metal oxides and their _

structural transformations: (A1) well-ordered R3m structure, (A2) the cation disorder or cation _

_

mixing phase with Fm3m structure, (A3) R3m structure with Li vacancies at highly charged states, (A4) partially cation mixed phase with transition metal ions in Li slabs. Adapted with permission [45]. Copyright 2015, Wiley-VCH. (B) Triangular phase diagram of LiNiO2–LiCoO2–LiMnO2 and the corresponding compositions. Adapted with permission [53]. Copyright 2017, The Royal Society of Chemistry. (C) Correlation among discharge capacity, thermal stability and capacity retention of Li1–δNixCoyMnzO2 materials. Adapted with permission [54]. Copyright 2013, Elsevier.

The Li+ ions mainly diffuse along the two-dimensional interstitial space in the Li+

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slabs between the TmO2 slabs, due to the strong chemical bonding between the metal ions of Ni, Co, Mn and Al and O2– ions in the TmO2 slabs [44, 45, 55, 56]. Firstprinciples investigation of the Li diffusion within the layered LixTmO2 (0 ˂ x ≤ 1) shows that: (1) for most values of x, Li+ ions migrate to adjacent vacant octahedral sites through intermediate tetrahedral sites, which is mediated through a divacancy mechanism; (2) the activation barrier (i.e., the difference in energy between at the activated state and at the initial equilibrium state of the hopping) associated with the divacancy hopping mechanism depends strongly on the Li concentration resulting in a diffusion coefficient variation within several orders of magnitude [57-59]. Shaju et al. [60] observed that the Li diffusivity of the NCM333 materials was 3×10–10 cm2 s–1 at 25 oC and changed in the charge-discharge process at 0.1 C. Some other factors including the Ni content, thickness of the Li slab (usually 0.26–0.27 nm) and cation mixing also severely affect the Li diffusion process. 2.2. Roles and effects (structure/chemistry/property) of metal ions The main metal elements in the cathode materials have different electrochemical redox potentials: (1) Ni is the mostly important active element, which can be oxidized from Ni2+ to Ni3+ at a potential of 3.6–3.9 V (vs. Li/Li+) and then to Ni4+ at a potential of 3.9–4.4 V; (2) Co is also electrochemically active and can be oxidized from Co3+ to Co4+ at a potential of 4.6–4.8 V; (3) Mn4+ and Al3+ remain inactive during the electrochemical cycling [40]. The average oxidation states of the transition metal ions with the detailed Li concentration (also related to the voltage change) can also be calculated by DFT computations based on the projected density of state and spin density distribution at different delithiation states of the NCM/NCA materials , and there is also a small magnetic moment on the oxygen atoms [61-63]. There is a trade-off relationship among the capacity, cycling stability and thermal property in the layered cathode materials (Fig. 2B-C) [53, 54, 64]. A higher Ni content could offer a higher initial reversible capacity; however, Ni2+ ions with a similar radius of 0.69 Å to that of Li+ ions (0.76 Å) can easily diffuse into the Li slabs and be oxidized to smaller Ni3+/4+ ions during the charge process (i.e., Li/Ni cation mixing) [40]. Firstprinciples calculations suggest that more Li+/Ni2+ ion exchange will shrink the Li interslab space thickness, and cause higher Li+ ion migration barrier [65]. The Ni atoms are hard to move after the Li/Ni exchange, which is detrimental to Li+ transport and the cycling performance. Besides, the NCM cathode materials with high Ni content (>75% Ni) suffer from more serious Li/Ni mixing during the cycling [66]. The Ni4+ ions in the charged cathodes would be spontaneously reduced to stable Ni2+ ions, which leads to undesired side reactions with the electrolytes and formation of oxygen-containing compounds like O22–, O2–, O– and O2 [67-69]. The local structural distortion caused by oxygen defects should be the driving force, which may lower the energy barrier of Ni2+ and Li+ exchanging positions and enable the structural relaxation [70, 71]. The highly10

delithiated cathode materials are also subjected to exothermic and endothermic phase _

_

transition from the layered R3m structure to spinel Fd3m structure or even rock-salt Fm _

3m structure, accompanied with the oxygen and carbon dioxide gas release (Fig. 2A) [50, 72-74]. Schipper et al. [75] calculated the diffusion barriers and paths of Ni2+ from octahedral (Oh) sites (in the Tm layer) to the corresponding tetrahedral (Td) sites (adjacent to the Li layer) by varying the Li vacancy number, and proved the rapid and irreversible Ni2+ ion migration upon lowering the Li ion concentration (Fig. 3). They also proposed a mechanism to explain the layered to spinel structure transition: a Ni2+ ion from an Oh in a Tm layer firstly migrates to a corresponding Td’’ site (step 1, Fig. 3B), followed by an Oh to Td migration of a second Ni2+ (in an adjacent layer) (step 2, Fig. 3C-D). Then, the second Ni2+ ion migrates to a corresponding Td’ site (Td–Td’) though a little higher barrier of ~0.7 eV (step 3, Fig. 3C-D). Finally, the Ni2+ ion would migrate from Td’ site to an Oh site (in the Li layer) by following cycling or via a Li concentration gradient, because of higher stability of Oh site than Td site with lowering in the amount of Li vacancies.

Fig. 3 Energy profiles for Ni migration obtained using PBE. Oh to Td Ni migration barrier in LiNi0.6Co0.2Mn0.2O2 with Li (A) di-vacancy near a Ni Td site and (B) tri-vacancy near the Ni Td site. (C) Oh–Td–Td’ Ni migration barrier with Li tri-vacancy with one Ni at the Td’’ site (above the migrating ion layer). (D) A proposed mechanism for Ni2+ migration causing a partial spinel phase. Adapted with permission [75]. Copyright 2016, Royal Society of Chemistry.

Apart from the charging condition, the exposure to the electrolyte for moderate amounts of time or aging in the absence of electrolytic solution without the electrochemical cycling would also cause the surface reconstruction of the cathode

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materials, which is proved by both the density functional theory (DFT) calculations and experimental characterization technologies [76]. The disordered rock-salt like structure needs higher activation energy for Li+ ion diffusion in comparison with the perfectly _

ordered R3m structure because of its smaller distance between the slabs, which makes Li+ ions more difficult to pass through [77]. When the Ni content increase, the exothermic decomposition temperature always shifts to lower temperature along with huge heat generation [54]. Particularly, the Ni-rich materials at highly delithiation state incline to suffer from oxygen evolution from the crystal lattice at temperatures ranging 150–300 oC [40]. The increase of the Ni content (≥50 mol%) brings about partial formation of Ni3+ from Ni2+ to meet the charge neutrality, which is confirmed by the proof that almost all of the Ni2+ ions in LiNi0.4Co0.3Mn0.3O2 change to Ni3+ ions in LiNi0.8Co0.1Mn0.1O2 and LiNi0.8Co0.15Al0.05O2 [40, 44, 78]. Based on the first-principles calculations, Liang et al. found that the valence of Tm ions in the ternary materials was determined by the atomic environment, the Tm-Tm bond strength was in the order of Mn4+Mn4+ > Ni2+Mn4+ > Ni3+Mn4+ > Co3+Mn4+ > Co2+Mn4+ > Ni2+Ni4+, and the formation of Co and Mn clusters in Ni-rich materials (Ni ≥ 80 mol%) was responsible for the phase degradation during cycling [79]. The incorporation of smaller Co3+ ions (0.545 Å in radius) could reduce the Li/Ni cation mixing, suppress the phase transition and increase the electrical conductivity, therefore improving the structural stability, cycling stability and rate performance at the expense of significantly increased material cost [42, 50]. The partial incorporation (5 mol%) of electrochemically-inactive Al3+ (0.50 Å in radius) can improve the safety by enhancing the structural and thermal stability particularly at highly-delithiated sates [44, 50]. Grey et al. [80] reported that the Ni3+ ions induced a dynamic Jahn-Teller distortion and the Al can reduce the strain resulted from the distortion by preferential electronic ordering of the Jahn-Teller lengthened bonds directed toward the Al3+ ions. However, high content of the inactive Al3+ in the TmO2 slabs is detrimental to the expansion/shrinkage of the lattice structure upon Li intercalation/deintercalation processes and leads to the lower capacity [81]. Mn exists as Mn4+, one of the most stable transition metal cations in close-packed oxygen frameworks, because of the filling of only the lower manifold of d-states [82]. The inactive Mn4+ ions (0.53 Å in radius) can reduce the transition metal dissolution and the irreversible side reactions between the electrode surface and electrolyte and thus provide significant structural stability and safety without the occurrence of the JahnTeller effect associated with Mn3+ ions [43]. Large amount of Mn4+ ions would lead to a high ratio of Ni2+/Ni3+, promotion of the Li/Ni cation mixing, and the phase transition _

_

from R3m to F3m during the electrochemical cycles [45, 83]. When decreasing the Mn content in the total amount of transition metal ions, the operation voltage in the initial

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charging stage decreases, but the discharge voltage are similar for all the compositions [84]. Recent comparison study of LiNi0.8Co0.15Al0.05O2 and LiNi0.7Co0.15Mn0.15O2 showed that Mn substitution appeared far less effective than Al in inhibiting the active element dissolution and irreversible phase transitions of the layered cathode materials [85]. In addition to the Ni2+ ion disorder in the Li slabs, Doeff et al. [34] recently studied the surface-deterioration mechanism of LiNi0.4Mn0.4Co0.18Ti0.02O2 during the cycling process by using electron energy loss spectroscopy (EELS) analysis, and found the cation mixing layer consisting of not only Ni2+ ions but also of other metal ions such as Mn2+ and Co2+. Hoang et al. [86] studied the intrinsic point defects in LiNi1/3Co1/3Mn1/3O2 and LiNi1/3Co1/3Al1/3O2 by hybrid density functional computations. In LiNi1/3Co1/3Mn1/3O2, the Ni and Li antisites (i.e., NiLi+ and LiNi‒) are low-energy defects apart from the Tm antisites on the Tm sublattice, and the formation energies follow the order of NiLi+ < CoLi+ < MnLi+. In LiNi1/3Co1/3Al1/3O2, possible low-energy antisite defects are NiLi0, AlLi0 and CoLi0 in addition to the Tm and Al antisites on the Tm sublattice, and the formation energies are in the order NiLi0 < AlLi0 < CoLi0. The low energy of Ni (Co)/Li antisites can be partially attributed to the similar ionic size of Ni2+ (Co2+) to Li+, while the low energy of the Al antisite is ascribed to the strong and directional interaction of the highly delocalized Al 3p states across the Li and transitionmetal oxide layers. X-ray diffraction (XRD) and neutron diffraction are usually used to detect the degree of the cation disorder [87]. The disordering leads to a partial destructive interference of the (003) plane’s constructive interference at a Bragg angle of θd(003) and a decrease in intensity of the (003) peak (Fig. 2A). However, the intensity of (104) peak increases since the transition metal ions in the Li layer are also on the (104) plane, resulting in increased constructive interference of the peaks for the (104) planes. Consequently, the intensity ratio of the (003)/(104) peaks (i.e., I(003)/I(104)) decreases when the degree of the disordering increases [45]. Generally, when both the ratio of the lattice parameters (i.e., c/a) and the I(003)/I(104) value are higher than 4.9 and 1.2, respectively, the NCM materials show a low cation mixing degree [38]. Furthermore, the crystallinity of the NCM materials is reflected by the split of the diffraction peaks of (006)/(102) and (108)/(110), and the higher splitting degree results in higher crystallinity and better cycling performance [88, 89]. The cation migration from the transition metal sites to Li sites occurs not only during the material synthesis procedure but also the electrochemical cycling process. Maintaining a proper Li+/Ni2+ ratio in the NCM synthesis process could effectively decrease the cation mixing degree and increase the crystallinity, and thereby improve the electrochemical properties. Currently the partial cation disordering is reconsidered as an appropriate way to stabilize the surface of the active materials based on the following reasons [45]: (1) the low Li+ ion content in the cation mixed phase causes a higher chemical stability and

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reduces the unstable side reactions with the electrolyte; (2) the Ni2+ ions in the Li sites act as “pillars” and provide an electrostatic repulsion force to hinder the further migration of the transition metal ions to the Li sites. 2.3. Effects of the preparation process Solid phase sintering, co-precipitation, sol-gel and spray-drying are usually utilized to synthesize the layered NCM/NCA cathode materials, and these methods possess various characteristics [90]: (1) the solid phase sintering route based on physical mixing and high-temperature crystallization for long time is simple for large-scale production of microsized particles, but it is difficult to control the elemental homogeneity, size distribution and morphology of the particles [91, 92]; (2) the two-step process including the metal precursor synthesis by co-precipitation and following lithiation with lithium salts at high temperatures of ≥700 oC is widely used to synthesize the secondary particles with homogeneous size distributions, highly controlled stoichiometric compositions and porous structures [36, 93-97]; (3) the sol-gel method is conducive to achieving uniform distribution of the ions and sub-microsized/nanosized particles with narrow size distribution [98, 99]; (4) spray-drying is also a useful process for synthesizing homogeneous spherical precursors with the metal element mixing at atomic level, and the subsequent calcination of the precursors can also result in singlephase cathode particles with fine distribution and uniform particle morphology [100102]. The preparation conditions can affect not only the stoichiometric ratio and distribution of the transition metal elements, but also other physical properties such as the morphology, size distribution, specific surface area and tap density of the layered cathode materials [103-114]. For instant, the pre-oxidation of the hydroxide precursors by Na2S2O8 in the co-precipitation process can increase the average valence of the transition metals (Co and Ni), decrease the Li+/Ni2+ mixing degree, and induce a highlyordered crystal structure [115]. The utilization of novel synthesis methods such as ion exchange and hydrothermal assisted process are also reported to effectively inhibit the Li/Ni mixing. As another example, synthesizing smaller or nanostructured particles can decrease the Li-ion diffusion paths and alleviate the inner stress caused by the volume change during lithiation/delithiation, and thereby enhance the rate performance and cycling stability. Note that the Li salts of Li2CO3 or LiOH in the preparation process are excessive to compensate the Li source loss during the calcination step and minimize the cation mixing for well-defined crystal structure [116]. The residual Li salts (Li2O) react with CO2 and H2O during the storage period in air atmosphere to form Li compounds such as LiOH and Li2CO3 [117, 118], which not only severely affect the slurry coating process due to the composite gelation caused by the high pH values of the NCM solutions, but also result in the increase of the electrode impedance [119, 120]. The residual Li compounds also promote the side reactions with the electrolyte solvents and

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acidic species generated from the decomposition of the polymer binders and are responsible for the undesirable gas evolution of CO2 during the fully charged state or at high temperatures [121-125]. Although the elimination of these Li residues by washing with water or surface coating is required for high performance [126], the ionic exchange between Li+ and H+ should be minimized during the removal process of the Li residues. Additionally, the washing process not only increases the processing time and cost, but also makes the cathode materials more chemically sensitive to moisture and air and less thermally stable than the unwashed samples [126, 127]. 2.4. Other changes (side reactions and morphology) during cycling process Apart from the surface contaminations from the Li residues in the material synthesis and storage processes, spontaneous side reactions also occur on the cathode surfaces during the electrochemical cycling, and the formed solid-state electrolyte interface (SEI) passive layers composed of various organic and inorganic electrolyte decomposition products are detrimental to the Li+ ion diffusivity during cycling and the battery performance (Fig. 4A) [45, 120, 128-131]. The by-products of oxidation at the cathode can also migrate to the anode surface and be reduced there, leading to the consumption of Li+ ions, reduction in Li inventory, and thickening of SEI on the anode [123, 132135]. Zheng et al. [136] made quantitative analyses of the electrode surface evolution for a commercial 18650-type LiNi0.5Co0.2Mn0.3O2/graphite cell during 3000 cycles under normal operation, and found that the Li inventory loss at the cathode accounted for ~60% of the cell capacity loss at different cycling stages and the missing transferable Li is immobilized on the graphite surface due to the SEI growth. Actually, the formation of surface passive layers on the cathodes similar to that generated at high voltage operation also occurred even when the cathodes are immersed in the electrolyte [44]. In addition, the highly reactive Ni4+ and Co4+ at high voltage of over 4.5 V could catalyze the decomposition of electrolyte which is accompanied by the structure/phase transformation and oxygen generation, leading to the electrolyte venting, smoke, fire or exploration especially at high temperature of over 55 oC [137, 138]. The Ni migration from the bulk to the surface and subsequent reactions with the electrolyte at a faster rate than Mn will also lead to voltage fading, owing to the decreased redox ability and a reduction in stabilizing Mn-Ni interactions [139]. The first-principles calculations of the atomic charges by Min et al. showed that the oxygen atom in cathode materials with more Ni content was already more oxidized initially and its variation from initial to full delithiation became larger (to satisfy the charge neutrality), indicating the increased possibility of the oxygen gas evolution in Ni-rich layered structures [140]. Streich et al. [141] developed an online electrochemical mass spectrometry to measure the gas release from NCMs of varying Ni content. They found that there were almost no gas evolution until 4.3 V regardless of the Ni content in NCMs, but both the CO2 and O2 gas were substantial between 4.3‒4.7 V and also correlated with the Ni content. All CO2

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and O2 evolution onsets fell into a very narrow SNOX (i.e., state of Ni oxidation) range of 85‒100% (Fig. 4B). Moreover, both χ (rise coefficient) and rO2,lim (upper limit) increased linearly with the Ni/Co ratio of NCM. A phenomenological model based on first-principles calculations was further used to describe the evolution of the valence band structure of NCMs during charging (Fig. 4C). The highest energy bands of NCMs included antibonding Ni eg and bonding Co t2g, Ni t2g, Mn t2g and O 2p states. During the charging, these bands were sequentially becoming depleted of electrons, progressively moving the Fermi level (EF) toward lower energies. The reactive oxygen formation triggered by Mn t2g or Mn−O2* oxidation was of minor importance, while the reactive oxygen formation was faster based on Ni eg and Ni−O2* than based on Co t2g and Co−O2*. Thus, the rate of reactive oxygen formation and r(CO2) and r(O2) was most controlled by the rate of electron depletion from the Ni-O2* surface states, and then was attenuated by the competitive oxidation of the Co t2g bulk and the Co-O2* surface states to a smaller extent. The density functional theory (DFT) calculations furtherly prove that the oxygen exposure on the electrode surface can greatly lower the Fermi level to the level of electrolyte and therefore accelerate the decomposition of electrolyte on the electrode surface [142].

Fig. 4 (A) The composition and structure change of the layered cathode materials during cycling. Adapted with permission [45]. Copyright 2015, Wiley-VCH. (B) Dependence of CO2 (top) and O2 (bottom) evolution on the state of Ni oxidation (SNOX) during 1st charge. The dashed gray lines indicate the SNOX range into which the gas evolution onsets fall. (C) Qualitative electronic density of state (DOS) diagrams for NCM111 (left of ordinates) and NCM811 (right of ordinates) at full lithiation, maximum Ni oxidation, and complete delithiation. The dashed horizontal line indicates

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the Fermi level (EF), and the occupied and unoccupied states are represented by filled and hollow ovals, respectively. Adapted with permission [141]. Copyright 2017, American Chemical Society.

Besides, the electrode materials are subjected to repeated anisotropic distention and shrinkage during electrochemical cycling under various depth of discharge and temperatures [125, 143-146], and the inner strain/stress may cause the formation of voids/pores in/between the primary particles (Fig. 5) [17, 147-151]. Generally, during the charge process, the c-axis increases because of the increase of the repulsive force between the oxygen slabs, while the a-axis decreases initially because of the increase of the electrostatic attraction force resulted from the oxidation of the transition metals. The increase in Ni content in the cathode materials will also result in more severe evolution of the c-axis, due to the extraction of a large amount of the Li ions [152-154]. The firstprinciples computations on the spin density distribution of NCM422 also show that the sharp absolute stress increase during the delithiation period is most probably ascribed to the local magnetic environment change associated with strong spin-flip transition of Ni ions caused by large Li+/Ni2+ exchange [61]. Furthermore, the electrolyte would easily penetrate into the pores and the dissolution of the active materials near the grain boundaries could lead to the collapse and isolation of the primary particles, resulting in the degradation of the electrical properties and the formation of Ni-O like impurity layers [16, 148, 155, 156]. Miller et al. [157] in situ observed that micro-crack and inter-grain separation even occurred during the first delithiation process at high C rate and the electrolyte permeated inside the LiNi0.76Co0.14Al0.10O2 particles through the crack, which resulted in the increased polarization and capacity degradation. Cho et al. [158] furtherly found that the collapse of each primary particle induced the thickness change of the overall electrode (~31% expansion after 300 cycles at 60 oC for ~2.7 g cm–3 electrode density).

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Fig. 5 A schematic model for the micro-crack growth and deterioration of a LiNi0.76Co0.14Al0.10O2 secondary particle during cycling. Adapted with permission [17]. Copyright 2014 Elsevier.

2.5. Performance degradation mechanisms By scrutinizing the crystal structure/composition/morphology/properties of the NCM/NCA ternary cathode materials and changes mainly in the materials synthesis and battery cycling processes, the performance degradation mechanisms of the NCM/NCAbased cathodes can be summarized as follows: (1) In the materials synthesis and storage processes (Fig. 6A), the similar radius size of Li+ and Ni2+ ions induces the Li/Ni exchange in the NCM/NCA lattice structure, which would shrink the Li interslab space thickness and cause higher Li+ ion migration barrier. Meanwhile, the Li residuals such as LiOH and Li2CO3 generate with the exposure of excessive Li salt raw materials in environmental conditions, which cannot only hinder the Li+ ion transport, but also facilitate the side reactions with the electrolyte, polymer binder and conductive additive and the generation of gas; (2) In the cathode preparation process (Fig. 6B), the mechanical impacts on the NCM/NCA secondary microparticles in the slurry stirring and electrode pressing stages would cause the particle rupture due to the poor interactions between the primary particles, resulting in the loss of the electrical contact in the cathodes and much more side reactions with the electrolyte during the battery charge/discharge processes; (3) In the cell fabrication process (Fig. 6C), the exposure of the NCM/NCA materials to the electrolytes for moderate amounts of time can lead to the side reactions with the electrolyte on the cathode surface in the absence of the electrochemical cycling, the dissolution of metal ions by the electrolyte attack, and even the surface reconstruction of the NCM/NCA materials;

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(4) The main changes occur in the battery cycling period (Fig. 6D). During the charging (delithiation) process, Li+ ions move to the anode side through the separator/electrolyte. The Li deficiency in the cathode materials would easily induce the cation mixing, oxygen loss and phase change. At the same time (usually 3.6‒4.8 V vs. Li/Li+), the oxidization of Ni2+/3+ and Co3+ to higher states (Ni3+/4+ and Co4+) and also the spontaneous reduction of the high-valence ions would easily catalyze the electrolyte decomposition to generate by-products and SEI films on the cathode surface, and the formation of oxygen-containing compounds such as O22–, O2–, O– and O2. The oxygen defects in the crystal structure would lower the energy barrier of Tm (Ni/Co/Mn/Al) and _

Li exchanging positions, and thus result in the phase transition from the layered R3m _

_

structure to spinel Fd3m structure or even rock-salt Fm3m structure, accompanied with the oxygen and carbon dioxide gas release. The gas evolution may lead to the electrolyte venting, smoke, fire or exploration especially at high temperature of ≥55 oC. The by-products of oxidation (e.g., Tm- and Li-based organic compounds) at the cathode can migrate to the anode surface and be reduced there, and result in the consumption of Li+ ions and thickening of SEI on the anode. The Li inventory loss at the cathode materials accounts for most of the cell capacity loss at the electrochemical cycling stage. Besides, the varied repulsion/attraction forces of the oxygen layers and Tm ions during the continuous lithiation/delithiation processes would induce the changes of the lattice parameters. The repeated anisotropic distention and shrinkage and the resulted inner strain/stress would cause the formation of voids/pores in/between the primary NCM/NCA particles and the rupture of the secondary particles, weakening the electrical contact in the cathodes. The electrolyte can easily penetrate into the voids/pores and then corrode the NCM/NCA materials, and furtherly accelerate the adverse physical/chemical phenomena. In short, the changes (or instability/irreversibility) of the cathode materials composition (e.g., Li deficiency/depletion, metal dissolution and oxygen evolution), structure (e.g., cation mixing, phase transition and SEI formation) and morphology (e.g., grain micro-crack and particle rupture), and the corresponding undesirable physical/chemical processes (e.g., side reactions and electric contact loss) hinder the charge transfer (mainly for Li+ ions) and cause the poor performance of the NCM/NCA ternary cathode materials.

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Fig. 6 Performance degradation mechanisms of the NCM/NCA cathodes during the (A) materials synthesis/storage, (B) electrode preparation, (C) cell fabrication, and (D) battery cycling processes.

2.6. Remarks To briefly summarize, there is a trade-off relationship between the capacity, cycling stability and thermal performance in the layered NCM/NCA cathode materials. Particularly, the Ni-rich materials show more promising commercial application prospect in terms of theoretical capacity, working voltage and synthesis cost, however, more severe reactions and changes in the surface and bulk structure and even in the electrolyte happen during the synthesis, storage and charge-discharge processes, leading to poor cycling/thermal stability. These problems are in fact resulted from the instability of the materials composition, structure and morphology, and the consequent undesirable

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physical/chemical phenomena. To inhibit these detrimental changes, in addition to the strict control in the materials synthesis and storage, electrode preparation and cell assembly processes, further designs at the NCM/NCA materials, electrode, and electrolyte levels are becoming more and more necessary (Fig. 1). 3. NCM/NCA material design As we mentioned above, the large-scale commercial application of the NCM/NCA cathode materials is still limited by the cation mixing, oxygen release, transition metal dissolution, surface side reactions, low cycling stability especially at high voltage or temperature, low conductivity and rate performance, poor compatibility with electrolyte, low tap/packing density, etc [159, 160]. Apart from the strict regulation of the synthesis condition, the NCM/NCA materials need to be further engineered by body doping, surface coating, compositing, etc. on the composition/structure/morphology levels to greatly improve their structure stability for better performance (Fig. 1 and 7) [161-163].

Fig. 7 NCM/NCA material design methods.

3.1. Composition design–body doping with other elements Atomic doping or elemental substitution is considered as a simple and important strategy to improve the structure stability of the NCM/NCA cathode materials [164]. Commonly, there are three kinds of sites for substitution: lithium sites in lithium ion layers, transition metal sites and oxygen sites in the transition metal layers [165]. The elemental substitution usually occurs in the precursor synthesis process and the hightemperature sintering process. In addition to the experimental studies, the firstprinciples computations also play an important role in investigating the effects of the elemental doping on the physical and electrochemical properties such as crystal

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structure, redox potentials, ion mobility, possible phase transfer and structural stability changes [166]. More and more researchers are using this method to optimize or minimize the experimental preparation and testing, and to predict the material performance under diverse conditions. However, it should be noted that the calculated results should be vindicated by the experiments, because it is still difficult to completely describe the complicated experimental conditions by first-principles calculations and the limited computational capabilities of modern computers [167]. There is also an ongoing search for exchange-correction functions to accurately describe the electronic structures and predict the material properties. 3.1.1. Cation substitution The structural and thermal stabilities and cycling performance can be greatly improved by doping the cathode materials with a few metal elements such as Al, Mg, Ti, Zr, Ru, Cr, Ga, Yb, Fe and Mo, which have similar ion radius to the Ni2+/Co3+/Mn4+ ions. The doping can reduce the content of the instable elements such as Li and Ni, prohibit the Ni transport from the transition metal slabs into the Li slabs by stabilizing the valance state of Ni and the electrostatic repulsion reaction, prevent the unwanted phase transition from the layered structure to spinel or rock-salt structure, enhance the bond strength between the transition metal ions and oxygen ions, or promote the Li+ ion transport owing to the enlarged Li slab distance [45]. It is essential to strictly control the element type, doping content and doping method, which severely determine the properties of the doped cathode materials. Partial substitution of Al3+ (0.50 Å in radius) has been proven to be an effect and low-cost method in modifying the electronic structure and improving the electrochemical performance of the cathode materials [168]. The Al doping cannot also reduce the react activity between the cathode materials and electrolyte, but also enhance the structural stability and inhibit the cationic migration and oxygen evolution due to the strong binding energy between Al and O [169]. From a thermodynamic viewpoint, the Gibbs energy for the formation of Al2O3 at 25 oC (–1582.3 kJ mol–1) is sufficiently to stabilize the crystal structure. Moreover, the Al doping is very advantageous, because the atomic weight of Al is lighter than that of other transition metal ions. The Al substitution at the transition metal sites is not easy because of the slower phase formation reaction rate of Al than these of Li, Ni and Co and the easy formation of undesired phases of LiAlO2 and Al2O3 during the conventional solid-state reaction, so solution-based synthesis methods such as co-precipitation and sol-gel are preferred to obtain pure NCM or NCA phase [16]. Liu et al. [170] investigated the effects of Al and Fe doping on LiNi1/3Mn1/3Co(1/3–x)MxO2 (M=Fe or Al, and x=0, 1/20, 1/9 or 1/6) synthesized by firing the co-precipitates of metal hydroxides and the metal oxides dopants. The lattice parameter of c value of the Al-substituted NCM increased while the a value slightly decreased with the increasing Al content, because of the similar ionic

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radius of Al3+ and Co3+; low-content (x<1/20) Al doping also improved the capacity retention of the LiNi1/3Co1/3Mn1/3O2 electrode by suppressing the cycling-induced lattice strain. Different from the Al doping, when increasing the Fe content, lattice parameters of both a and c increased due to the larger Fe3+ radius, and the cycling stability became poor because of the occupation of Fe3+ in 3a sites, block of Li+ transport and increase of electrode polarization. Zhou et al. [171] also reported the effect of Al doping on LiNi1/3Mn1/3Co(1/3–x)AlxO2. When adding very little Al, the NCM material charged to 4.3 V reacted dramatically with 1 M LiPF6 EC/DEC electrolyte when the temperature exceeded 260 oC; when the doping content x increased from 0.04 to 0.1, the self-heating rate decreased from 20 to 3 oC min–1, indicating the Al substitution for Co in the NCM can improve the thermal stability. However, the specific gravimetric capacity and volumetric energy density of the doped NCM (140 mAh g–1 and 2.42 Wh cm–3) were lower than the undoped (163 mAh g–1 and 2.87 Wh cm–3). Recently, Aurbach et al. [172] studied the effects of Al doping on LiNi0.5Co0.2Mn0.3O2. The Al substitution in the Ni sites in the host structure can improve the crystallinity, suppress the oxygen evolution from the host structure and inhibit the Ni migration from the transition metal sites to the Li sites. Consequently, the Al-doped electrode displayed a lower capacity fading of 0.02% per cycle than the undoped electrode (0.07% per cycle). When cycled at 45 oC, the both the electrodes showed comparable capacities at moderate rates, while the Aldoped electrode delivered slightly superior capacities at higher currents. The Al-doped electrode also showed lower charge transfer resistance after the electrochemical cycling, owing to the enhanced structure stability and reduced side reactions with the electrolyte. The study from the Manthiram group furtherly disclosed that substituting 2 mol% Al for Ni in LiNi0.92Co0.06Al0.02O2 can suppress the transition metal dissolution and prevent the side reactions for better battery performance [173]. Boron ions (B3+) have the same valence to Al3+ but lower weight, which could reduce the adverse impact from the incorporation of inactive atoms to some extent. As we know in Section 2.2, the transition metal (Tm) ions can migrate from the octahedral sites in the Tm layers to the tetrahedral sites in the Li layers, and furtherly migrate to the octahedral sites in the Li layers, which would hinder the Li diffusion. Because of the smaller radius of B3+ ions (0.27 Å), they can be intercalated into the tetrahedral sites in the Tm layers and block the migration of Tm ions from the Tm layers to the Li layers, and hence enhance the structural stability. Sun et al. [174] found that the B doping did not change the intrinsic crystal structure of Li-rich Li1.2Ni0.13Co0.13Mn0.54O2. However, the calculated energy barriers of Mn ion migration in B-doped sample was six times of that for the B-free sample, indicating that the B substitution greatly block the Mn migration process. Because of the enhance structural stability by the B doping, the Bdoped electrodes showed high initial reversible capacity of 293.9 mAh g‒1 and capacity retention of 89.5% after 100 cycles at 0.5 C within 2.0‒4.6 V.

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Mg2+ has a similar radius of 0.72 Å to Li+ and is also an important dopant for stabilizing the crystal structure of the cathode materials [175-181]. The first-principles calculations show that Mg induces stronger and less distorted Ni-O bonds than Al does and the stronger Ni-O bonds hinders the collapse of the original crystal structure to the NiO-like phase during delithiation; however, the calculated atomic arrangement for Mgdoped materials predicts higher activation energy for Li migration and implies that the Mg doping could make the structure kinetically more persistent against delithiation [179, 182, 183]. Li et al. [179] found that the Mg doping in LiNi0.8Co0.15Al0.05O2 resulted in the increase of the lattice parameters of a and c, the higher activation energy for the Li migration, and the suppression of the phase transition between H2 and H3. Kim’s research about the Mg-doped LiNi1/3Co1/3Mn1/3O2 found [184]: (1) the parameters of a and c values slightly decreased for the Ni site substitution, while the a and c values slightly increased for the Co and Mn site substitution; (2) the Mg substitution for Ni resulted in relatively lower tap density, while the tap density increased when the Mg substitution for Co and Mn; (3) partial Mg substitution for Ni and Co resulted in decreased capacity and relatively poor cyclability, while the substitution for Mn significantly improved the obtained capacity, cyclability, rate capability and thermal stability at highly oxidized state. Recent studies on the Mg-doped cathode materials suggest that Mg is prone to occupy the Li sites and would restrain the Li/Ni cation mixing and the phase transition [45, 185, 186]. A Mg-doped Li1.2Ni0.12Co0.12Mn0.56O2 sample showed a lower initial specific gravimetric capacity than the undoped one, because of the following factors: (1) the Mg2+ occupied the Li sites instead of the transition metal sites; (2) the electrochemically inactive Mg2+ did not participate in redox reaction; (3) compared to Li-O bond, the formation of stronger Mg-O bond can suppress the extraction of Li+ and O2 from Li2MnO3 component during the initial charge process. Nevertheless, the Mg-doped electrode of Li1.2Ni0.12Co0.12Mn0.536Mg0.024O2 exhibited a higher capacity of 236.9 mAh g−1 than the undoped electrode (187.0 mAh g−1) after 30 cycles at 12.5 mA g−1 in a voltage of 2.0−4.8 V, owing to the stabilization of the layered structure by the Mg2+ doping. In order to systematically investigate the effects of Al and Mg doping on the formation of oxygen vacancies and cation disordering, and the structural stability of Nirich cathode materials, Min et al. [187] obtained the formation energies based on the first-principles calculations. It was found that the Al and Mg substitutions at the Ni and Li sites were the most preferred (with the lowest formation energies), respectively (Fig. 8A), which also agreed with the reports [45, 172, 185, 186]. Oxygen vacancies can be more effectively prevented in an Al-doped structure while cation disordering (i.e., NiLi and LiNi–NiLi defects) can be effectively suppressed by Mg-doping in a fully lithiated NCM. The Mg-doped structure exhibits the smallest decrease in the total volume up to full delithiation (‒9.61 Å3, ‒1.03%) than the Al-doped (‒10.97 Å3, ‒1.18%) and

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undoped (‒14.56 Å3, ‒1.56%) structures from the initial stage of delithiation up to around 66%. The Mg dopant in the Li layer can also act as a supporting pillar to mitigate the collapse of the Li layer during delithiation. Based on the charge analysis of oxygen atoms (Fig. 8B) and formation energy calculations (Fig. 8C-E), they also concluded that the oxygen stability improved the most in the Al-doped structure regardless of whether the structure had pre-existing defects or was delithiated, while the Mg doping can effectively inhibit the formation of excess Ni and Li/Ni exchange under the same defective conditions. Moreover, the delithiation process would cause the possibility of formation of all types of defects especially the oxygen vacancies, and any type of defect can induce the formation of the others. Dixit et al. [188] furtherly found that the Li diffusion barrier near the Al-doped sites (0.63 eV) was higher than that of undoped NCM523 material (0.50 eV), due to the greater repulsive interactions between Li+-Al3+ ions. However, they thought that the increase in Li-diffusion barrier on doping was a local effect, because the c lattice parameter did not change significantly. They also found that the Li diffusion barrier in the vicinity of the second-nearest neighboring Tm (Ni, 0.53 eV) was only slightly higher than that of the undoped material, and thus concluded that the low Al doping level (5 at%) was unlikely to affect the bulk Li diffusion property.

Fig. 8 (A) The formation energy of Al and Mg doping when doped at Ni, Co, Mn, and Li sites (the insets are the atomic configurations of Al and Mg substitutions at Ni and Li sites, respectively), (B)

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the oxygen charge variation at each layer during delithiation for undoped, Al-doped and Mg-doped structures (inset is the supercell structure of NCM doped with Al and Mg), and the formation energies of (C) oxygen vacancy, (D) Ni migration and (E) Ni migration after the generation of oxygen vacancy for the fully lithiated and partially delithiated (66% and 69% delithiation for the Al and Mg-doped, respectively) structures. Adapted with permission [187]. Copyright 2017 Royal society of chemistry.

A few other divalent metal ions (e.g., Cu2+, Zn2+ and Ca2+) also play a similar role to Mg2+ as cationic dopants in affecting the electrochemical performance of the layered cathode materials [189-191]. Chen et al. [192] synthesized a series of Ca-doped LiNi0.8– 0.8xCo0.1Mn0.1Ca0.8xO2 (x=0–8 mol%) by a traditional solid-state method. The doping of Ca2+ (radius of 0.1 nm than Ni2+) can decrease the Li/Ni disorder of the cathode materials, and increase the Li+ diffusion coefficient compared to the pristine LiNi0.8Co0.1Mn0.1O2. As a consequence, the Ca-doped sample with an optimized composition (x=6 mol%) showed higher capacity of 122 mAh g−1 and capacity retention of 81% than the un-doped one (91 mAh g−1 and 69%) after 30 cycles at 0.2 C in a voltage of 3.0−4.3 V. Moreover, the Ca doping can evidently improve the cycling stability even at a higher voltage of 4.5 V, lower the electrode polarization, and alleviate the voltage reduction during cycling. Zhao et al. [193] studied the structure, morphology and electrochemical performance of pristine and Zn-doped Li1.2Ni0.13–x/3Co0.13–x/3Mn0.54– 2+ (0. 74 Å) to Li+, the lattice x/3ZnxO2. Because of the similar ionic radius of Zn parameter of c increased after the Zn doping, which indicated the enlarged interlayer spacing and would facilitate the Li intercalation/deintercalation and prevent the structure deterioration. The morphology of the samples did not changed with the increasing Zn content with the secondary microparticles composed of 300–400 nmsized primary particles. The initial specific capacity of the samples decreased with the increasing Zn content, owing to the two factors: (1) the inactive Zn2+ would occupy the Li sites and did not participate in redox reaction; (2) the stronger Zn-O bond can suppress the extraction of Li+ and O2– from Li2MnO3 during initial delithiation process. However, due to the stabilization of the layered structure and restraint of the polarization during cycling, the Zn-doped electrode (x=0.07) displayed higher capacity of 137.3 mAh g−1 and capacity retention of 83.6% than the pristine one (101.6 mAh g−1 and 78.6%) after 100 cycles at 1 C in a voltage range of 2.5–4.8 V. Aliovalent substitution with Ti4+ (0.605 Å in radius) and Mn4+ is more complicated than the corresponding isovalent substitution processes (e.g., with Al3+ and Fe3+) owing to the need for charge compensation. Kim et al. [194] suggested that the introduced Ti4+ ions rendered the structure rigid because of the sufficiently low Gibbs energy of –888.8 kJ mol–1 for the formation of TiO2 at 25 oC. Kam et al. [195] think the influence of the aliovalent substitution of Co3+ with Ti4+ on the electrochemical properties of

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Li[Ni1/3Co1/3−yTiyMn1/3]O2: (1) Li co-substitutes with Ti on transition metal sites to compensate for charge and results in Li-excess materials, and the higher discharge capacity after charging to 4.7 V is probably attributable to an activation process similar to that seen in composite NCM materials containing a Li2MnO3 component, in which both Li and oxygen are removed at high potentials; (2) the charge compensation occurs by a different mechanism such as the replacement of a small amount of Mn4+ with Mn3+, and the higher capacity is not contingent upon any activation process. To more clearly explain the enhanced cycling performance of Li[Ni1/3Co1/3−yTiyMn1/3]O2, Markus et al. [196] applied the DFT calculations including Hubbard-U corrections. It was found that Ti substitution resulted in a smaller decrease in the overall volume change during the cell cycling (Fig. 9A). The larger Ti4+ substitution of Co3+ also led to the formation of a charge-compensating electron polaron on a nearby Mn3+ (Fig. 9B-C), expanding the surrounding TmO6 octahedrons and reducing the structural distortions during delithiation. Because of the extra electrons provided by the Ti substitution, the quasiequilibrium Li intercalation voltage also decreased (Fig. 9E), reducing the overpotential required to drive Li deintercalation and allowing more Li to be accessed for increased capacity (Fig. 9D). Moreover, the Ti substitution can help to stabilize the NCM structure by elevating the formation energy of the rock salt structure (Fig. 9F) and more strongly binding oxygen. Li et al. [197, 198] synthesized Mn-doped LiNi0.8Co0.15Al0.05O2 particles through an in situ oxidizing-coating method using KMnO4 as both Mn source and oxidant, and the Mn doping can reduce the Ni3+ to Ni2+ that can suppress the capacity fading and increase the thermal stability. Nurpeissova et al. [199] prepared Ti-doped LiNi0.8Co0.15Al0.05O2 microparticles (LiNi0.8Co0.15Al0.02Ti0.03O2, NCAT) by co-precipitation followed by sintering. The introduction of Ti did not alter the crystal structure but enhanced the structural integrity; furthermore, the NCAT electrode showed a lower occupation of Ni2+ in Li layers of 2.1% than the NCA electrode (4.2%) after 50 cycles at 1 C in a voltage range of 2.5–4.5 V. The Ti4+ ions acted as charge compensator and structural stabilizer, which can maintain the valence states of the transition metals at around 3+, hinder Ni2+ diffusion into Li sites, and prohibit the formation of surface layers of NiO phase (an electronic/ionic insulator). As a result, the NCAT electrode showed higher initial capacity of 203 mAh g–1 and capacity retention of 74% than the NCA electrode (192 mAh g–1 and 67%) after 50 cycles at 1 C. The NCAT electrode was also more suitable for higher capacity delivery at high currents than the NCA electrode.

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Fig. 9 (A) Change in the Li interstitial volume as a function of Li concentration for unsubstituted (Ti00) and Ti-substituted (Ti03: 3 mol%) NCM, (B) localized polaron by plotting the total density of states (DOS) for unsubstituted and Ti-substituted NCM, including the projected DOS for the Mn3+ cation, (C) the charge density difference map between unsubstituted and Ti-substituted NCM (the colored bands represent charge depletion for the blue bands and charge accumulation for yellow and orange bands), (D) experimental quasi-equilibrium voltage for unsubstituted and Ti-substituted NCM by integrating data from a stepped potential experiment, (E) voltage difference between the intercalation voltage for unsubstituted and Ti-substituted NCM as calculated using DFT and experimentally using the stepped potential experiment, and (F) calculated free energy of formation at room temperature for the rock salt structure from NCM structures with different Li concentrations. Adapted with permission [196]. Copyright 2014 American Chemical Society.

The radius of Zr4+ is 0.72 Å similar to the Li+ radius (0.76 Å) but larger than the Ni2+/Co3+/Mn4+ radius, and the bond energy of Zr-O is stronger than that of Ni/Co/MnO, and thus Zr4+ can easily enter the Li slabs and function as “pillar” to enhance the crystal structure [200-203]. Schipper et al. [75] found that Zr4+ ions preferred to substitute in Li and Ni sites by a combined experimental and first-principles computational study. The number of Ni3+ ions decreased and the number of Ni2+ ions increased upon the Zr substitution at Ni sites, owing to the compensation of the excessive positive charge from the aliovalent Zr4+ dopant. The band gap of LiNi0.6Co0.2Mn0.2O2 with the Zr doping also disappeared due to the broadening of the Ni and Co states. Furthermore, the Zr-doped cathode retained the layered structure upon cycling at 45 oC. However, the Zr substitution for Mn in NCM shows no significant improvement in thermal stability [204]. The excessive doping of Zr could also result in the formation of impurity phases (e.g., LiZrO3) and therefore cannot improve the

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cycling performance [200, 201]. Ding et al. [205] prepared LiNi1/3Co1/3Mn1/3-xZrxO2 (x=0, 0.01, 0.025, 0.05) via a thermal polymerization method. All the samples have a hexagonal α-NaFeO2 structure and the lattice parameters increased monotonously with the Zr content. The Zr doping can not only stabilize the lattice structure, but also lead to higher Li diffusion coefficient in NCM. Consequently, the LiNi1/3Co1/3Mn1/3-0.01Zr0.01O2 showed the best electrochemical performance with a capacity retention of 92.7% after 100 cycles and a capacity of 133.9 mAh g−1 at 8 C, corresponding to 71.5% of its capacity at 0.1 C. Wang et al. [206] discussed the effects of Zr content on the structure and electrochemical properties of Li(Ni0.5Co0.2Mn0.3)1–xZrxO2 powders. When increasing the Zr doping content, the lattice parameters of a and c increased. Moreover, the thickness of the Li slabs increased while the thickness of transition metal slabs decreased with the increase of the Zr content, which would facilitate the Li+ ion transport and improve the structural stability. Therefore, the Zr-doped electrodes with x=0.005 and 0.01 displayed higher capacities of 145 and 155 mAh g−1 and capacity retentions of 77% and 84% than the un-doped electrode (130 mAh g−1 and 69%) after 100 cycles at 1 C in a voltage range of 3.0−4.6 V. When x=0.01, the Zr-doped electrode also exhibited much more enhanced rate capability among all the samples. Contrast to the comparatively integrated morphology of the Zr-doped sample, the un-doped LiNi0.5Co0.2Mn0.3O2 particles suffer severe damage and collapse into small fragments after 100 cycles under 4.6 V. Vanadium (V) doping can greatly improve the rate capacity of the layered cathode materials because of the increase of Li+ ion diffusion coefficient [207]. A V-doped LiNi0.5Co0.2Mn0.3O2 cathode at 5 C remained 72.2% of the capacity at 0.1 C, while the capacity retention was only 58.9% for the undoped electrode; the V-doped electrode also exhibited a higher capacity retention of 96.3% after 50 cycles at 1 C than the undoped one, because of the hindrance of Ni/Li mixing by the V substitution in 3b sites [208]. Lu et al. [209] prepared V-doped Li-rich Li1.2Mn0.52–x/3Co0.08–x/3Ni0.2–x/3VxO2 cathode materials by a solid state thermal decomposition method using generated watersoluble vanadyl oxalate as dopant. The V-doped sample (x=0.015) showed higher capacity of 152.2 mAh g−1 and capacity retention of 90.2% than the undoped one (140.2 mAh g−1 and 88.3%) after 50 cycles at 1 C in a voltage range of 2.0–4.8 V. The doped electrode also showed a higher average capacity of 99.0 mAh g−1 than the undoped electrode (66.2 mAh g−1) at 5 C. The performance improvement for the doped sample was ascribed to the following reasons: (1) vanadium ions with larger radius (0.59 Å) could enlarge the interplanar spacing and facilitate Li+ ion insertion/extraction inside the layered structure; (2) the much stronger V-O bonds than Ni-O, Co-O and Mn-O bonds can enhance the structure stability; (3) the vanadium ions doped into the crystal lattices possess two valence states of +4 and +5, and the fast electron transfer between V4+ and V5+ can enhance the electronic conductivity and reduce the electrochemical polarization.

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The Cho group [210] furtherly found that the heat treatment of LixNi0.8Co0.15Al0.05O2 with ammonium vanadate precursors can not only reduce the Li residuals on the NCA surface but also result in the formation of 17 nm thick surface layer with V4+ ions in the transition metal sites for decreasing the thickness of the Li/Ni mixing layer. Therefore, the V-treated NCA electrode demonstrated excellent capacity of 179 mAh g−1 after 200 cycles between 3.0−4.3 V with a capacity retention of 90%, which was 18% higher than that of the untreated NCA electrode. High-valence metal ions such as Mo6+ (0.62 Å in radius) and Se6+ (0.42 Å in radius) have also been reported as dopants to significantly affect the structure of the cathode materials for improving the electrochemical properties [95, 211-215]. It is known that the d-bands of the transition metals are pinned at the top of the oxygen p-bands in the transition metal oxides, which affects the structural stability of the oxides by the evolution of O2 gas. Doped metals with a high valence of 6+ can strongly hybridize with the oxygen ions and reduce the hybridization between the Ni and O atoms, and therefore enhance the structural stability and impede the O2 evolution [216]. A series of Se6+-doped Li1.2[Mn0.7Ni0.2Co0.1]0.8–xSexO2 nanoparticles were prepared by coprecipitation and subsequent calcination [217]. Because the bonding energy of Se-O (465 kJ mol–1) is higher than those of Mn-O (402 kJ mol–1), Ni-O (392 kJ mol–1) and Co-O (368 kJ mol–1), the Se-doping can enhance the structural stability of the layered cathode materials, suppress the oxidation process of O2– to O2, and hinder the layeredto-spinel phase transformation to some extent. Consequently, the Se-doped electrodes maintained larger first Coulombic efficiency (77%), more stable performance, higher rate capacity (178 mAh g−1 at 10 C) and higher mid-point voltage retention (95% after 100 cycles) in comparison with the pristine electrode. The layered cathode materials are also doped with a few rare earth elements such as La, Ce and Pr [218]. Ding et al. [219] synthesized a series of cathode materials with a formula of Li[Ni1/3Co1/3Mn1/3]1–xRexO2 (Re=La, Ce or Pr, x≤0.04) by a sol-gel method using citric acid as chelating agent. The lattice parameter ratio of c/3a increased slightly as the La content increased, indicating the distortion of lattice of LiNi1/3Co1/3Mn1/3O2. When the La content increased from 0 to 0.04, the lattice parameters of a and c increased from 2.862 to 2.876 Å and from 14.184 to 14.262 Å, respectively, which would facilitate the extraction and insertion of Li+ ions from the active materials. Therefore, the doped samples (Re=La, Ce or Pr, x=0.03) showed higher initial capacity of more than 160 mAh g–1 and capacity retention of 97% than the pristine sample (154 mAh g–1 and 95%) after 20 cycles at 1 C in a voltage range of 2.6–4.4 V. Yu et al. [220] prepared La-doped Li-rich Li1.2Mn0.54–xNi0.13Co0.13LaxO2 (x=0.01, 0.02, 0.03) by a solvothermal method and subsequent calcination technique (Fig. 10A-E). The Mn4+ sites were effectively substituted with La3+ ions, and thus can expand the pathway for intercalation and deintercalation of Li+ ions. The La doping can also stabilize the

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layered framework and suppress the layered-to-spinel transformation upon long term cycling. Consequently, the doped electrode delivered a high specific capacity of 286.4 mAh g–1 at 0.1 C with a high Coulombic efficiency of 90.4% and excellent cycling performance with 93.2% capacity retention, compared to 58.4% for the undoped electrode after 100 cycles at 1 C in a voltage range of 2.0–4.6 V (Fig. 10F). Even when cycled at 5 C after 100 cycles, the doped electrode showed higher capacity of 150.0 mAh g–1 and capacity retention of 88.8% than the undoped electrode (75.9 mAh g–1 and 54.2%). Furthermore, the average voltage of the undoped electrode decreased to 2.72 V after 100 cycles at 1 C, while the voltage fading of the doped electrode was mitigated and the average voltage only decreased to 3.13 V (Fig. 10G). The Raman characterization of the La-doped sample with less band shift than the undoped sample further proved that the La doping indeed mitigated the layered-to-spinel phase transformation (Fig. 10H).

Fig. 10 (A) XRD patterns, (B) SEM images, (C) TEM images, (D) EDS spectra, (E) FFT patterns of the selected area marked in red rectangle in (C), (F) cycling performance and (G) the corresponding average discharge voltage at 1 C between 2.0–4.6 V, and (H) Raman spectra of the bare and Ladoped Li1.2Mn0.54–xNi0.13Co0.13LaxO2 (x=0.01, 0.02 and 0.03, denoted as L0-LLO, L1-LLO and L2LLO, respectively). Adapted with permission [220]. Copyright 2016 Elsevier.

In addition to a part of the aforesaid metal ions, a few alkali metal ions (e.g., Na+ and K+) can also be introduced into the Li layers to enlarge the Li slab space for enhancing the structural stability of the layered cathode materials and facilitating the Li+

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ion diffusion [221-226]. Chen et al. [227] investigated the effect of Na doping on the Li1–xNaxNi1/3Co1/3Mn1/3O2. When x=0.03, all the Na+ ions (1.02 Å in radius) occupied the Li sites, reduced the degree of the cation mixing, and stabilized the crystal structure, and thus the doped electrode displayed a higher capacity of 141.8 mAh g–1 than the undoped electrode (132.9 mAh g–1) after 50 cycles at 0.5 C in a voltage range of 2.5–4.4 V. Hu et al. [228] synthesized Na-doped LiNi0.8Co0.15Al0.05O2 by co-precipitation and following solid state calcination using Na2CO3 as a Na resource. The Na ions occupied the Li slabs and did not change the particle morphology. Although the Na doping resulted in a little decrease in the initial discharge capacity, the Na-doped electrodes showed improved capacity retention and superior rate performance. The Li0.99Na0.01Ni0.8Co0.15Al0.05O2 electrode displayed an initial capacity of 184.6 mAh g–1 at 0.1 C but higher capacity retention of 90.71% than the pristine one after 200 cycles at 1 C, because of the enlarged Li layer spacing and the decreased cation mixing degree. Ates et al. [229] found that the larger crystal lattice of Na-doped Li-rich NCM most likely provided better transition metal ion migration (most probably Ni3+) into the Na depleted sites (due to the exchange between the larger Na+ in the NCM and Li+ in electrolyte) in the Li layers after the activation of Li2MnO3 in the first charge process and thereby stabilized the layered structure against the undesirable conversion to the spinel phase and improved the cycle life and rate capacity. Soon afterwards, Li et al. [230] scrutinized the effect of K+ doping on the Li-rich Li1.20Ni0.13Co0.13Mn0.54O2, which was prepared by a solid state sintering method using K+-doped α-MnO2 as a starting material. They found that the doped potassium did not give rise to the variation of the grain size and chemical valence state of the pristine materials, but the interslab thickness of LiO2 increased slightly after the K+ doping and thus would reduce the activation barrier for Li hopping and facilitate Li+ ion migration. More importantly, the K+ doping can effectively mitigate the transition from the layered structure to the spinel variant, because of the following reasons: (1) the doped K+ ions acted as fixed pillars in the Li layers and might weaken the formation of tri-vacancies in the tetrahedron sites (i.e., 8a sites) in the Li layers and Mn3+ ion migration from octahedron sites (16d sites) to the tetrahedron sites; (2) the larger K+ ions (0.77 Å in radius) can possibly aggravate steric hindrance for the spinel growth based on the kinetic consideration of crystallography. As a consequence, the doped electrode exhibited higher capacity of 283 mAh g–1 and capacity retention of 91% than the undoped electrode (245 mAh g–1 and 81%) after 30 cycles at 20 mA g–1 in a voltage range of 2.0–4.8 V. Besides, the doped electrode also exhibited slower voltage decline and better rate capability than the K+-free electrode. As we have mentioned that a few ions such as Zr4+ (0.72 Å) with the radius similar to the Li+ radius can also easily enter the Li slabs and function as “pillar” to enhance the crystal structure [200-203]. Feng et al. [231] investigated the impact of Ti4+ substitution

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in the Li-layers of the Li1.2Ni0.13Co0.13Mn0.54O2 on the structure and cycling performance by DFT calculations and materials characterizations. The Ti doping cannot prevent the Mn migration and the layered to spinel-like transformation during long-term cycling; however, the Ti doping can increase the LiO2 interslab spacing and Li ion diffusion rate, and reinforce the structural stability of the formed spinel phase during the cycling. As a result, the 2.5% Ti-doped electrode showed higher capacity of 229 mAh g‒1 (71% retention) after 300 cycles at 0.2 C and rate capacity of 136 mAh g‒1 at 5 C than the undoped electrode (81 mAh g‒1 at 5 C) within 2.0‒4.8 V. 3.1.2. Anion substitution Apart from the cation doping, the doping with some anions (e.g., F–, Cl–, Br– and S2–) can also improve the cycling performance of the cathode materials [161, 232-237]. The cation doping site is difficult to control since several transition metals with different oxidation states coexist in the transition metal layer; however, the anions can be highly incorporated into the bulk structure without a significant sacrifice in reversible capacity since there is only one kind of anion of O2– in the cathode materials [238]. Among them, sulfur substitution of oxygen was found to induce a more flexible structure in the layered cathode materials, because S2– has lower electronegativity than O2– and the relatively large size and polarizability of S2– make it easy for Li+ ions to transport in the layered structure and reduce the structural strain during lithiation/delithiation [239]. Especially, the effect of F– substitution on the cathode materials has been well studied: (1) the substitution of the O2– ions (1.40 Å in radius) with the smaller F– ions (1.33 Å in radius) will decrease the average valence state of the transition metal ions and meanwhile increase the average radius of the transition metal ions; (2) the stronger bond character of Li and F may result in repulsive force in the oxide matrix, so the lattice would expand both in a and c axes and thus enhance the structural and thermal stability during the lithiation/delithiation [240-242]; (3) the fluorine substitution can catalyze the growth of the primary particles due to the effect of LiF on crystallization, which in turn results in high tap density as well as high volumetric capacity [238, 243-245]. Guo et al. [246] found the fluorine doping in the NCA crystal lattice can decrease the cell parameters and interlayer spacing because of the small radius of F– and the large electronegativity. The entire bond energy and structure stability can be greatly enhanced by the partial replacement of the metal-oxygen bond with the metal-fluorine bond in the NCA materials [247]. Wang et al. [248] thought the fluorine substitution in oxygen sites can lead to the formation of the cubic rock structure (NiO-like phase) on the cathode surface, reduce the release of oxygen, and suppress the side reactions between the cathode and electrolyte. Meanwhile, the fluorine ions could react with Li salt residues on the cathode surface during the doping process, and therefore would inhibit the formation of HF and restrain the side reactions between the cathode materials and electrolyte. Moreover, the fluorine doping could increase the lattice parameters and

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promote the intercalation/deintercalation processes of the Li+ ions. Therefore, the doped electrode of LiNi0.73Co0.12Mn0.15O2–xFx (x=0.02) showed higher capacity of 155 mAh g– 1 and capacity retention of 97.5% than the undoped electrode (140 mAh g–1 and 87.4%) after 200 cycles at 1 C in a voltage range of 3.0–4.3 V. As we have mentioned in Section 2.2, a few researchers have reported the positive effects of Ni/Li mixing by hindering the further structural degradation though the inhibition of Li migration by the antisites. Li et al. [249] scrutinized the structure and properties of halogen-doped LiNi0.85Co0.075Mn0.075O2 and furtherly verified the importance of the antisites in enhancing the cathode stability by combined study of firstprinciples calculations and neutron powder diffraction. The halogen (especially F) substitution can promote the neighboring Li and Ni atoms to exchange their sites, leading to the formation of a more stable halide-based local octahedron of [(Ni2Li1)halogen-(Li2Ni1)] (Fig. 11A). In this octahedron model, a halogen atom formed three bonds with 2 Ni atoms and 1 Li atom in the Ni layer, and meanwhile formed three bonds with 2 Li atoms and 1 Ni atom in the Li layer. There was a linear relationship between the F doping content and antisite concentration in NCM. Moderate F substitution (1%) induced 5.7% antisites in the NCM, but excessive F doping can cause the generation of irreversible rock-salt phase. Besides, Mn/Li and Co/Li antisites were not found in the LiNi0.85Co0.075Mn0.075O2-xFx sample. The F doping resulted in the decrease of the electronic conductivity; however, the lattice parameter of c increased with the F doping (Fig. 11B) and would offer a wider path for Li+ ion migration. Compared with the pristine sample, the calculated Mulliken charges of O in local layered structure (below approximately −0.8 e per atom) of the F-doped NCM material became more negative (Fig. 11C), which resulted in the increased electrostatic repulsion force between oxygen layers and then the lattice parameter of c. The Ni atoms near F would also get less positive to keep the valence equilibrium after the O2− substitution with F−, and then lead to the mixing valence of cations, which can facilitate the Li+ transfer. Due to the improved structure stability and Li+ ion diffusivity, the 1% F-doped NCM electrode showed high capacity of 153.4 mAh g−1 and capacity retention of 83.3% after 120 cycles at 1 C and 55 oC within 2.8‒4.3 V (Fig. 11D). SEM characterizations also proved the high morphology/structural stability of the F-doped samples (Fig. 11D insets).

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Fig. 11 (A) The effect of F, Cl, Br and I substitution on the calculated antisite defect formation energies with the local octahedron structure models, (B) neutron Rietveld refinement and the corresponding ground state crystal structure of the 1% F-doped NCM, (C) the Mulliken charges of oxygen in pristine and F-doped NCMs, and (D) the cycling performance of pristine and 1% F-doped NCM cathodes at 1 C and 55 oC between 2.8‒4.3 V (insets: SEM images of the NCMs after the cycling). Adapted with permission [249]. Copyright 2018 Wiley-VCH.

Polyanion-type compounds are a class of materials in which tetrahedral polyanion structure units (XO4)n− and their derivatives (XmO3m+1)n− (X = B, P, Si, S, As, Mo, or W) with strong covalent bonding combine with MOx (M=transition metal) polyhedrons [250-252]. Therefore, these polyanions are used as dopants to improve the cycling performance and thermal stability of the layered cathode materials, due to the merits of the polyanions for stabilizing the TM(3d)−O(2p) bonds and oxygen close-packed structure [253, 254]. Qu et al. [252] synthesized boracic polyanion-doped NCA materials of [LiNi0.8Co0.15Al0.05](BO3)x(BO4)yO2-3x-4y using H3BO3 as boron source by pre-coating treatment and subsequent solid calcination process. The boracic polyanions were incorporated into the bulk structure and more enriched in the surface layer. Since the effective suppression of the particle destruction and SEI formation, the NCA sample with x+y=0.015 showed high capacity of 155.1 mAh g–1 and retention rate of 96.7% after 200 cycles at 2C. Cong et al. [255] investigated the structure and properties of the (PO4)3−-doped layered NCM materials (i.e., LiNi1/3Co1/3Mn1/3O1.94(PO4)0.015) prepared by co-precipitation and subsequent sintering. The tetrahedral (PO4)3−-doped sample with enlarged lattice parameter of c can provide broader Li+ ion transport paths and thereby

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improve the rate capability (e.g., 98 mAh g–1 at 3000 mA g–1 between 2.8–4.3 V). Moreover, the (PO4)3− ions with robust P-O bonds in the host structure can suppress the oxygen evolution, cationic rearrangements and crystalline phase transition, and thus could ensure the long-term cycling stability. For example, the (PO4)3−-doped electrode exhibited higher capacity of 112 mAh g–1 and capacity retention of 80% than the undoped electrode (88 mAh g–1 and 63%) after 600 cycles at 300 mA g–1 in a voltage range of 2.8–4.5 V. Lately, Zhang et al. [256] prepared SiO44–- and SO42–-doped Li-rich cathode materials (denoted as LNCMO) by spray-drying and following solid-state reaction. The substitution of the larger polyanions (240 pm for SiO44– and 258 pm for SO42– in thermochemical radius) for the small O2– anions (140 pm in radius) in the R3m layered structure could result in lower cation mixing of Li+/Ni2+, better layered structure and a few crystal defects, which would be beneficial for blocking the structure transition during lithiation/delithiation. As a result, the polyanion-doped oxides had higher initial Coulombic efficiency and improved cycling stability. More specifically, the LNCMO(SiO4)0.05 and LNCMO-(SO4)0.03 electrodes showed higher capacities of 220.4 and 215.4 mAh g–1 and energy densities of 650 and 600 Wh kg–1than the undoped electrode (159.9 mAh g–1 and 450 Wh kg–1) after 400 cycles at 30 mA g–1 in a voltage range of 2.0–4.8 V, respectively. The potential decline of the LNCMO-(SiO4)0.05 and LNCMO(SO4)0.03 samples (0.70 and 0.57 V) were much lower than that of the undoped electrode (0.81 V). The LNCMO-(SO4)0.03 sample also exhibited higher rate capabilities than the LNCMO-(SiO4)0.05 and undoped samples. Furthermore, the onset and peak temperatures of the polyanion-doped samples were higher than those of the undoped LNCMO sample, indicating the relatively strong bonding energy of the metal cations with polyanions and the enhanced thermal stability by the polyanion doping. 3.1.3. Co-substitution Apart from the mono-doping with either single cation or anion, doping the cathode materials with more than two kind of cations (e.g., Na-Nb, Al-Mo, Mg-Al and Cu-Ti) [257-260] or both the cations and anions (e.g., Mg–F, Al–F and Cr–F) can provide a synergistic effect in improving the electrochemical performance [261]. Kim et al. [262] synthesized Mg–F co-doped LiNi1/3Co1/3Mn(1/3–x)MgxO2–yFy (x=0–0.04, y=0–0.04) via co-precipitation and subsequent high-temperature heat-treatment. The Mg–F cosubstitution can enhance the crystallinity, increase the lattice parameters, reduce the cation mixing and increase the tap-density (2.43 g cm–3), and thus the Mg–F co-doped electrode displayed better cycling stability (180 mAh g–1 after 30 cycles at 20 mA g–1 between 2.8–4.6 V) and thermal stability than the undoped and mono-doped electrodes. Liao et al. [263] further investigated the effect of Al–F co-doping on the LiNi0.333Co0.333Mn0.293Al0.04O2–zFz particles synthesized by a sol-gel route and following sintering process: (1) although the radii of F– anions (1.33 Å ) is smaller than that of O2– (1.40 Å ), the substitution of fluorine resulted in the partial reduction of the transition

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metal ions for the charge compensation of F anions and the increase of the lattice parameters; (2) the particle size increased from 2 to 8 μm with the increase of the fluorine content; (3) the strong bond of Li–F could perturb the Li+ ion intercalation/deintercalation and decrease the initial specific capacity, but would improve the Coulombic efficiency and cycling stability; (4) the Al–F co-doping can suppress the decomposition of the electrolyte and enhance the battery safety. Consequently, the co-doped electrode of LiNi0.333Co0.333Mn0.293Al0.04O1.95F0.05 showed higher initial Coulombic efficiency of 92.4%, capacity of 150 mAh g–1 and capacity retention of 94.9% than the Al-mono-doped electrode (90.1%, 143 mAh g–1 and 84.6%) after 20 cycles at 0.1 C in a voltage of 3.0–4.3 V. 3.1.4. Remarks Body doping is a simple and important method for improving the structure stability and electrochemical/thermal properties of the cathode materials. The layered cathode materials are usually doped with a few common metal elements (e.g., Al and Mg) and even transition metal (e.g., Ti and Zr) and rare earth (e.g., La and Ce) elements mainly for substitution on the transition metal (Ni/Co/Mn) sites for improving the crystal structure stability. The Li sites in the layered cathode materials can be also substituted with a few alkali (and alkali earth) metal elements (e.g., Na, K and Mg) to enlarge the Li slab distance for facilitating the Li transfer and enhancing the lattice stability. However, the doping conditions must be carefully controlled because: (1) a few doping elements are not really involved in the crystal structure but in the form of impurity phases; (2) the doping elements may occupy more than one metal site in the crystal structure; (3) the high content of the exotic elements can even reduce the total capacity of the cathodes. Compared to the cation substitution, the anions are easier to be incorporated into the bulk structure without significant sacrifice in capacity. The anion substitution (e.g., F– and S2–) on the oxygen sites can affect both the crystal structure and morphology characteristics. Moreover, the substitution from single-ions to polyanions (e.g., SO42– and SiO44–) seems to better stabilize the lattice structure of the NCM/NCA materials. In addition to the mono-doping with either single cation or anion, doping with more than two kind of cations (e.g., Na-Nb, Al-Mo, Mg-Al and Cu-Ti) or both the cations and anions (e.g., Mg–F, Al–F and Cr–F) can greatly enhancing the structure stability by cosubstitution on the metal and oxygen sites. The future research will be also focused on the co-substitution, and the simultaneous substitution on the Li, transition metal and oxygen sites may shed another light on improving the electrochemical and thermal properties. 3.2. Structural design–surface coating and compositing with other materials Different from the body structure, the interface between the electrode materials and the electrolyte plays an vital impact on the electron transfer, Li+ ion diffusion, metal

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dissolution, phase transition, side reactions, electrolyte decomposition, etc [264, 265]. Appropriate surface coating with substances as layers on the electrode surfaces would protect the cathode particles against side reactions with the electrolyte, avoid the loss of transition metal ions (e.g., Co4+ and Mn4+) or of oxygen, hinder the oxygen evolution, suppress the phase transition, or simply improve the electronic and ionic conductivities of the powders [47]. This route has turned out to be more efficient than the doping process to improve the electrochemical properties and the safety of Li-ion batteries. A few inert substances such as oxides, fluorides and phosphates are commonly utilized as coating layers for cathodes particles. In general, the coating effect on specific cathodes is highly dependent on the type of coating substance, coating content, coating thickness, coating uniformity and heat treatment condition [266, 267]. In particular, an ideal coating layer should be uniformly covered with proper coating thickness, because too thick or uneven coating layers will impede the electronic and ionic conductivities. The surface treatment often consists of the exposure of the active materials in treatment solutions/atmospheres followed by annealing steps, which may result in composition or structure changes on the surfaces [40]. As for the coating route, researchers have developed various approaches such as mechanical mixing, sol-gel, pulsed layer deposition (PLD), chemical vapor deposition (CVD), radio frequency sputtering (RFS) and atomic layer deposition (ALD), which can also be classified into solid phase-, liquid phase- and gas phase-based routes [268-274]. Since the cathode particles (i.e., so-called secondary particles) are actually constituted with smaller primary particles, so the inner spaces in the secondary particles cannot be completely treated through the solid phase route, resulting in incomplete coating on the primary particles; however, the liquid phase and gas phase approaches can tune the cathode particles from the surface into the inside for complete coating on all the primary particles [210, 275-277]. Amongst these methods, the mechanical mixing of the cathode powders with the coating precursors and subsequent annealing of the mixture (i.e., dry coating process) is a simple and scalable process without the use of solvents, but it is difficult to build completely-covered coating layers on the cathode particle surfaces. An important method for uniform coating layers on the cathode particles is by a wet coating process (e.g., sol-gel or solution processing) including the dispersion of the cathode particles in the precursor-involved solution and the following drying and calcination of the mixture, but it is difficult to control the thickness of the coating layers. Notably, the relatively low temperature in the subsequent annealing process for treating the mixtures obtained by the dry/wet routes usually results in the weak bonding/adhesion between the coating layers and the cathode particles. Sputtering is a well-established physical deposition approach in the semiconductor industry, because of its wide range of coating materials and high coating efficiency; however, it may result in high non-uniformity of the coating layers and the aggregation of the cathode particles, due to the irregular

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exposure of the cathode particles to the sputtered materials [273]. Similar to the solution-based approach, CVD is also a feasible way to homogeneously modify the particle surface using mixture gas as source [278]. ALD is another effective technique for achieving ultrathin and highly-conformal coating layers of metal oxides and solid electrolytes with thickness control at atomic scale without blocking Li+ ion diffusion into the cathode particles, but it is impossible for scaling up to industrial grade at low cost [33, 279-281]. Other methods like melting impregnation and PLD are not wellcontrolled processes, and the resulted coating layers are lack of conformality, uniformity and completeness [282]. Thus, it is still essential to develop improved or novel coating routes to obtain complete, uniform, robust and controllable coating layers on the cathode particles at low cost and easy scale-up for the massive applications. 3.2.1. Oxides The typical oxides as coating substances are MgO, Al2O3, CuO, ZnO, TiO2, ZrO2, SnO2, CeO2, SiOx, Sm2O3, La2O3, V2O5, Cr2O3, Y2O3, Mo2O3, Co3O4 and Pr6O11, which can protect the cathode materials from direct contact with the electrolyte and therefore inhibit the corrosion by the electrolyte and the decomposition of the active materials [283-295]. As we know LiPF6 is the most widely-used electrolyte salt and can easily decompose to generate HF (PF5+H2O → POF3+2HF) in the presence of moisture. The generated HF readily attacks the cathode materials to promote gradual dissolution of the transition metal ions from the particle surfaces. The commonly-used oxide coating layers can scavenge the acidic HF and are transformed to metal fluoride layers (e.g., Al2O3+6HF→2AlF3+3H2O), thus reducing the acidity of the electrolyte and the deterioration of the cathode surface. The preservation of the cathode surface stabilizes the host structure of the cathode materials, allowing for stable charge-transfer resistance during successive charge-discharge cycles [296]. However, the Li+ ion diffusivity in these oxide coating materials is relatively low, and too thick coating layers cannot only decrease the cycling capacity but also be detrimental to the rate capability [221, 297299]. Amongst these oxides, SiO2 is abundant, cheap and environmentally friendly, and can also react with HF to protect the cathode particle from electrolyte attacking and enhance the structure stability during cycling. Chen et al. [300] prepared SiO2-coated LiNi0.915Co0.075Al0.01O2 particles by a simple one-step dry coating method in an airtight powder homogenizer. The SiO2 coating layer can effectively inhibit the side reactions between the NCA and the electrolyte and improve the structural stability. As a result, the 0.2 wt% SiO2-coated NCA electrode exhibited higher capacity of 181.3 mAh g–1 and retention rate of 90.7% than the pristine NCA electrode (178.1 mAh g–1 and 88.7%) after 50 cycles at 1 C. More importantly, the coated electrode delivered much higher capacity of 146.4 mAh g–1 and retention rate of 66.8% than the uncoated electrode (122.7 mAh g–1 and 54.3%) after 50 cycles at 1 C and 60 oC. 39

MgO coating on the cathode particles have been fabricated by depositing Mg precursors followed by calcination [301, 302], pulsed laser deposition [303] and sol-gel [99, 304]. During the coating process, a portion of Mg2+ ions would easily diffuse into the LiO2 layers of the layered cathodes, which results in the partial substitution of transition metal ions and stabilization of the crystal structure [305]. Meanwhile, the superficial MgO coating layer can decrease the activation energy of Li+ transfer reaction at the interface. Both the doping and coating effect are advantageous to improve the cathode performance. A Cr2O3-coated NCM cathode material was fabricated by Zhang et al. [306] via a sol-gel method followed by calcination, and the surface coating with 1.0 wt% Cr2O3 did not affect the NCM crystal structure (α-NaFeO2) of the cathode material compared with the pristine material. The Cr2O3-coated NCM cathode showed higher capacity of 140 mAh g–1 and capacity retention of 83.1% than the bare cathode (116 mAh g–1 and 72.5%) after 30 cycles at 0.5 C under a high cutoff voltage of 4.5 V, mainly because of the inhibition of the interface resistance between the cathode and electrolyte. Al2O3 coating layers of several nanometers in thickness have also been proven to effectively enhance the cycling performance [163, 307-309]. The Li1.2Ni0.13Mn0.54O2 particles coated with 1wt% Al2O3 and 1wt% RuO2 delivered 280 mAh g–1 at 0.05 C and 160 mAh g–1 at 5 C [310], because the addition of RuO2 may prevent the diffusion of Al into the particles and increase the metallic character of the surface layer, and the coating layer can resist the side reactions with the electrolyte. Hu et al. [311] synthesized Co3O4-coated LiNi0.8Co0.15Al0.05O2 particles with the coating thickness of 3–5 nm by a wet-chemical route and following annealing process. The Co3O4-coated NCA solution had a smaller pH value of 11.05 than the pristine one (11.38), because a few Li residuals of LiOH and Li2CO3 on the NCA surface reacted with Co3O4 during the annealing process. The partial elimination of the Li residuals can increase the conductivity and facilitate the Li diffusion. Additionally, the coating layer can suppress the dissolution of the transition metal from the NCA material. As a consequence, the Co3O4-coated NCA electrode displayed higher capacity of 175 mAh g–1 and retention rate of 91.6% than the pristine electrode (137 mAh g–1 and 79.1%) after 100 cycles at 1 C between 2.8–4.3 V. V2O5 is also an attractive coating material, because of its excellent electrical conductivity and high Li+ ion diffusion coefficient [285, 312-314]. V2O5 has been reported to improve the electrochemical cyclability of Li-rich layered cathode materials and LiCoO2 at high cut-off voltage, since the vanadium ions in 3d0 electronic states can reduce the surface catalytic activities and stabilize the surface oxide ions during their electrochemical oxidation [315, 316]. Surface coating of VOx can also suppress the Mn dissolution from spinel LiMn2O4 (60 ppm) [317]. Additionally, V2O5 coating layer could reduce the heat generation and increase the onset temperature of the exothermic

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reaction for the LiNi0.8Co0.15Al0.05O2, and thus showed high capacity of 179 mAh g–1 and capacity retention of 90% after 200 cycles at 60 oC [210]. He et al. prepared Li1.2Mn0.54Ni0.13Co0.13O2 particles coated with V2O5 layers by a wet chemical route [285]. The 3 wt% V2O5-coated sample exhibited higher capacity of 269.1 mAh g–1 and capacity retention of 96.3% than the uncoated sample (202.2 mAh g–1 and 80.2%) after 50 cycles at 25 Ah g–1 between 2.0 and 4.8 V, because of the effective reduction of the charge transfer resistance at the electrode-electrolyte interface and the improvement in the Li+ ion transportation in the particles. As well, TiO2 and ZrO2 coating are efficient for improving the cycling stability of the layered cathode materials during high potential cycling [318-323]. TiO2 coating did not affect the lattice of LiNi1/3Co1/3Mn1/3O2, but exhibited obvious effects on its discharge capacity and cycling stability [324, 325]. Zhou et al. [326] found that the sputtered TiO2 layers on LiCoO2 partially reacted with the decomposition product of electrolyte (e.g. HF) and formed a more stable and conductive interphase containing TiFx, which is responsible for the improvement of the cycle capability (160 mAh g–1 with 86.5% retention after 100 cycles at 1 C within 3.0–4.5 V). Qin et al. [33] furtherly coated ~5 nm-thick amorphous TiO2 layers on the LiNi0.6Co0.2Mn0.2O2 particles via an ALD approach (Fig. 12A-C). The TiO2-coated electrode showed higher capacity retentions of 85.9% and 80.8% with a capacity of 185.6 mAh g–1 than the uncoated electrode (67.5%, and 60.8% with a capacity of 107.8 mAh g–1) after 100 cycles at 1 C in a voltage range of 2.5–4.3 V at 25 and 55 oC, respectively (Fig. 12D-F). Such enhanced electrochemical performance of the coated sample was mainly attributed to the high-quality ultrathin coating of the amorphous TiO2, which can greatly protect the active material from the HF attack, suppress the dissolution of metal ions in the electrode and facilitate the Li+ ion diffusion through the electrode.

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Fig. 12 (A) Schematic illustration for TiO2 growth on the surfaces of NCM-622 particles by ALD, (B) SEM image with X-ray elemental area scanning and (C) TEM image of the TiO2-coated NCM particles, (D) rate capability, and cycling performance of the pristine and TiO2-coated samples at 1 C at various temperatures of (E) 25 and (F) 55 oC between 2.5–4.3 V, respectively. Adapted with permission [33]. Copyright 2016 The Royal Society of Chemistry.

Similar phenomenon also happened for the ZrO2-coated cathode materials [286]. Lee at al. [273] used a modified radio frequency sputtering (RFS) system for uniform coating of ZrOx layers on the LiNi1/3Co1/3Mn1/3O2 particles by continuous moving or rotation of the cathode powders during the sputtering process. When cycled at 1 C between 3.0–4.5 V, the uncoated electrode delivered an initial discharge capacity of 178 mAh g–1, whereas the ZrOx-coated electrodes delivered decreased capacities of 174, 168 and 159 mAh g–1 with the increasing sputtering time of 1, 2 and 3 hours, respectively, suggesting that the specific capacity at the first cycle was slightly sacrificed by introducing the ZrOx coating layers. However, the cycling stability was greatly improved by the ZrOx coating: the 1, 2 and 3 h-coated samples exhibited higher capacity retentions of 81.7%, 94.6% and 92.1% than the uncoated one (71.3%) after 70 cycles, respectively. Even when tested at 0.1 C between 1.9–3.7 V in all-solid-state cells with a configuration of LiNi1/3Co1/3Mn1/3O2/Li2S-P2S5/Ln2Li, the 1, 2 and 3 h-coated electrodes showed higher initial capacities of 86.4, 98.5 and 109.3 mAh g–1 and capacity retentions of 79.6%, 83.1% and 84.7% after 50 cycles than the uncoated one (77.4 mAh g–1 and 50.2%), respectively. The improved battery performance was attributed to the surface coating: physical/chemical protection of the active material surface, enhancement of the Li-ion diffusion kinetics and stabilization of the electrode/electrolyte interfaces. Coating the NCM/NCA cathode particles by metal oxides with high electrical conductivity and compatibility with the electrolyte are considered to greatly enhance the cycling stability and rate capability [327]. Du et al. [328] coated LiNi0.8Co0.15Al0.05O2 with Sb-doped SnO2 nanoparticles homogeneously and uniformly by a wet chemical method. The high-conductivity coating layers on the particles can not only impede the direct contact between the NCA particles and the electrolyte, but also offer high conductivity for facilitating the electron transport and improving the rate performance. Therefore, the SnO2-coated NCA particle with the coating thickness of 4–6 nm showed higher capacity of 165.4 mAh g–1 and retention rate of 91.7% than the uncoated particle (125.7 mAh g–1 and 70.9%) after 200 cycles at 1 C and 60 oC. Zhang et al. [329] completely and uniformly coated LiNi0.8Co0.15Al0.05O2 with nanoscale ZnO film by magnetron sputtering, and the ZnO-coated electrode showed high capacity of 169 mAh g–1 than the pristine electrode (127 mAh g–1) after 90 cycles at 1 C due to the reduced charge transfer resistance and effective protection of the NCA from cation dissolution.

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Apart from the commonly-used metal oxides, a few organic coating layers can also protect the electrode materials from side reactions in the electrolyte [330]. Wang et al. [331] coated an ethoxy-functional polysiloxane (EPS) on the LiNi0.6Co0.2Mn0.3O2 particle surface through the hydrolysis of tetraethyl orthosilicate (TEOS) and following polymerization during a wet ball-milling procedure (Fig. 13A and C-E). The hydrolysis-polymerization process can reduce the content of the residual Li salt and water on the cathode particle surface (Fig. 13B). Moreover, the EPS coating layer on the particle surface can react with HF in the electrolyte, suppress the accumulation of a thick SEI layer, and enhance the stability of the cathode surface. The 2 and 5 mol% EPS coated electrodes exhibited higher capacity retentions of 93% and 95% respectively than the uncoated electrode (87%) after 200 cycles at 1 C in a voltage range of 2.8–4.3 V (Fig. 13F). Even when cycled at 55 oC, the 1, 2 and 5 mol% coated electrodes showed higher capacity retentions of 81%, 85% and 93% respectively than the uncoated electrode (63%) after 100 cycles at 1 C (Fig. 13G). Additionally, the coated electrode showed a higher exothermic peak temperature of 310.90 °C and smaller exothermic peak than the uncoated electrode (284.97 °C), disclosing that the thermal stability of the NCM cathode material was greatly enhanced with the aid of the EPS coating, which can inhibit the interface reaction between the electrode and electrolyte and stabilize the interface (Fig. 13H).

Fig. 13 (A) Schematic illustration of EPS coating on the LiNi0.6Co0.2Mn0.3O2 particle surface (ENCM) through the hydrolysis of TEOS and following polymerization, (B) content of the residual Li

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and water on the NCM surface before and after the EPS coating treatment, (C) SEM/TEM images of 2 mol% E-NCM particles, (D) SEM of a E-NCM particle, (E) the corresponding EDS mapping in (D), cycling performance of the NCM samples at 1 C between 2.8–4.3 V at (F) 25 and (G) 55 oC, and (H) DSC profiles of the NCM and 2 mol% E-NCM cathodes at charged state to 4.3 V. Adapted with permission [331]. Copyright 2016 American Chemical Society.

3.2.2. Fluorides The metal oxide (e.g., ZnO and Al2O3) coatings can suppress the decomposition of electrolyte at the early cycling stage, scavenge the free HF species from the LiPF6 electrolyte decomposition to form metal fluorides, and suppress the dissolution of the transition metals. However, once the metal oxide coating layers react with HF, water is also generated and facilitates the HF generation. The generated HF would attack the cathode surface and this series of processes continues, further degrading the battery performance [163, 332, 333]. To tackle this problem, researchers attempt to directly use metal fluorides such as AlF3, MgF2, CaF2, LaF3, SrF2, YF3, CeF3 and FeF3 to coat the cathode materials [248, 332, 334-338]. For instant, AlF3 is well known to protect the Al cathode collectors from the corrosion by the LiPF6-based electrolyte, and the AlF3 coating can ensure the structural stability of the cathodes and promote the Li+ ion diffusion even at high operation temperature, due to the contribution of the rapid Li transferring role of the AlF3 layer [339-343]; LaF3 and CeF3 are more stable than most of the other compounds under the corrosion of HF, and they can significantly decrease the dissolution of the active cathode material [5, 344-347]; MgF2 coating could alleviate the dissolution of Co and the surface damage during cycling and enhance the thermal stability [345, 348-351]. Furthermore, these metal fluoride coatings could reduce the release of the oxygen from the cathode materials and inhibit the transformation of crystal structure from layer to olivine or even to rock salt phase, thus resulting to high thermal stability [352, 353]. Li et al. [354] synthesized 100–200 nm Li1.2Mn0.54Ni0.16Co0.08O2 particles via a sol-gel route and then coated them with 5–7 nm thick AlF3 layers through a wet-chemical process. The AlF3 coating can inhibit the formation of LiF films on the particles, suppress the increase of electrode-electrolyte interface resistance and reduce the release of oxygen, and thereby the AlF3-coated electrode displayed higher initial capacity of 208.2 mAh g–1 and capacity retention of 72.4% after 50 cycles at 1 C than the pristine electrode (191.7 mAh g–1 and 51.6%). Amine et al. [355] coated LiNi0.8Co0.15Al0.05O2 particles with 5 nm thick AlF3 (1 wt%) by a simple dry-coating process based on high-speed stirring. The AlF3-coated NCA/graphite full-cells showed higher capacity retentions of 86.2% after 1000 cycles at 25 oC and 55.9% after 500 cycles at 55 oC at 1 C than the pristine NCA-based full-cells (66.5% and 11.7%) between 3.0–4.2 V. Besides, the coated NCA showed higher exothermic peak of 237 oC and lower heat generation of 458 J g–1 than the pristine NCA

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(230 oC and 520 J g–1) charged at 4.3 V. The improvements in cycling and thermal performance were attribute to: (1) the AlF3 coating layer protected the NCA from HF attacking, reduced the metal dissolution, and suppressed the increase of the charge transfer resistance; and (2) the AlF3 layer inhibited the volume variation of the NCA upon lithiation/delithiation and prevented the particle pulverization. 3.2.3. Phosphates Relative to the aforesaid oxides and fluorides as coating substances, a few phosphates such as AlPO4, CoPO4, FePO4, Mn3(PO4)2 and Ni3(PO4)2 contain strong P=O bonds and can greatly suppress the corrosion of the phosphate-coated cathode materials from electrolyte; the strong covalency interactions between the polyanions of PO43– and metal ions can also enhance the thermal stability of the phosphate-coated cathode materials [138, 356-359]. Interestingly, a few early researches also showed that the metal phosphates can react with the Li residues at higher temperatures to generate olivine type Li conductive metal phosphates, which are very electrochemically and thermal stable even after full delithiation [103]. A few works have proven that the AlPO4 coating can better enhance the cycling performance of the cathode materials compared to the Al2O3 coating [138, 357, 360362]. Cho et al. [138] prepared AlPO4-coated LiNi0.8Co0.1Mn0.1O2 particles via coprecipitation and calcination. The DSC scans of the bare, Al2O3- and AlPO4-coated cathodes charged to 4.5 V showed that the bare cathode exhibited a steeply rising flow height up to 40 W g−1; however, the Al2O3- and AlPO4-coated cathodes produced approximately half and one quarter of the exothermic heat generated by the bare cathode respectively, indicating the AlPO4 coating layer can greatly minimize the exothermic reaction of the cathode with the electrolyte. Moreover, the AlPO4-coated electrode showed a higher capacity retention of ~100% than the bare electrode (92%) after 200 cycles at 1 C in a potential range of 3.0−4.2 V, because the AlPO4 coating can block the potential HF attack from the electrolyte and suppress the dissolution of metal ions (Ni, Co and Li) from the cathode surface. Ni3(PO4)2 is also found to effectively protect the cathode particles from HF attacking and suppress the increase of the charge transfer resistance. Scrosati et al. [363] prepared Ni3(PO4)2-coated LiNi0.8Co0.15Al0.05O2 particles with the coating layer of 20 nm in thickness by a simple dry-coating method based on ball milling. Because of the decrease of the Li/Ni cation mixing degree, inhibition of the dissolution of transition metals and enhancement of the crystal structure stability, the Ni3(PO4)2-coated NCA electrode showed higher capacity of 149 mAh g−1 and retention rate of 75% than the pristine NCA electrode (104 mAh g−1 and 53%) after 100 cycles at 0.5 C within 2.7−4.3 V at 55 oC. FePO4 is one of the environmentally-friendly, less expensive and thermally-stable materials, and has been coated on various cathode materials for achieving enhanced

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cycling performance especially at elevated temperature [364-368]. Liu et al. [369] fabricated FePO4-coated LiNi1/3Co1/3Mn1/3O2 particles with 20–30 nm thick FePO4 layers by co-precipitation followed by calcination, and also investigated the effect of preparation conditions (e.g., FePO4 content and calcination temperature) on the electrochemical property. The coated electrodes with 1 wt%, 2 wt%, 3 wt%, 4 wt% and 5 wt% FePO4 showed 138.4, 143.5, 132.1, 101.3 and 61.8 mAh g−1 after 100 cycles at 150 mA g−1 in a voltage range of 2.8−4.5 V, respectively, whilst the capacity was ~103.2 mAh g−1 at the 100th cycle for the pristine sample. The result indicated that only an appropriate amount of FePO4 coating with thin layers can facilitate the diffusion of Li+ ions and prevent from the side-reaction at the interface of electrode and electrolyte, while the excess insulating FePO4 with thicker layers may inhibit the intercalation/deintercalation of Li+ ions and lead to a worse electrochemical performance. The 2 wt% FePO4-coated electrodes calcinated at 300, 400, 500 and 600 oC showed 123.4, 143.4, 129.8 and 120.3 mAh g−1 after 100 cycles, respectively. The result can be explained as follows: with low heat treat temperature, the contact between the coating layer and the cathode material is not very close, so that the coating layer may peel off from the cathode surface during cycles; however, with high heat treated temperature, the coating layer may diffuse into the cathode, affect the crystal lattice and lead to bad Li+ ion transport between the electrode and electrolyte. LaPO4 coating can also suppress the loss of oxygen from the cathodes, and the absence of any transition element of the first series in the coating would prevent the dissolution of Ni, Co and Mn of the core region into the electrolyte, thus improving the cathode performance [47, 370, 371]. Park et al. [371] coated LiNi0.5Co0.2Mn0.3O2 powders with LaPO4 via co-precipitation and subsequent heat treatment. The 3 wt% LaPO4-coated electrode showed a higher capacity of 185 mAh g–1 after 60 cycles at 1 C within 3.0–4.8 V and better rate performance than the uncoated and 1 wt% and 5 wt% LaPO4-coated electrodes (150, 175 and 150 mAh g–1), due to the protective effect of the coating layer with regard to the reactive electrolyte. A few metal phosphates can also react with the Li residues at high temperatures to generate Li-containing equilibrium phases (as mentioned above), and the unreacted metal phosphates still can function as coating materials by preventing the direct exposure of the cathode material to the electrolyte. Min et al. [372] developed a firstprinciples-based screening process to identify ideal coating materials with higher reactivity to Li2O. Based on the reaction enthalpy values between 16 metal phosphates and Li, Co3(PO4)2, Mn3(PO4)2, Fe3(PO4)2, and TiPO4 were chosen for experimental validation. The experimental products generated after the reactions of the metal phosphates and Li were confirmed by comparison with predicted phases obtained from the phase diagram based on calculations. The Co3(PO4)2, Mn3(PO4)2 and Fe3(PO4)2coated LiNi0.91Co0.06Mn0.03O2 materials showed superior cycling performance after 50

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cycles than the TiPO4-coated materials, because of the better Li-removing capability of Co3(PO4)2, Mn3(PO4)2 and Fe3(PO4)2 than TiPO4. To better improve the electrochemical performance, double-layered shell structure on cathode materials are further developed. Wang et al. [373] prepared Li1.2Mn0.54Ni0.13Co0.13O2 particles with surface modified by double-layer coating with 2 wt% AlPO4 or CoPO4 inner layer and 2–3.5 wt% Al2O3 outer layer. The double-layer coated samples exhibited lower irreversible capacity loss (26 mAh g–1) and higher discharge capacity (295 mAh g–1) than both the uncoated (75 and 253 mAh g–1) and single-layer coated (27–49 and 269–285 mAh g–1) samples, ascribed to the retention of a higher number of oxide ion vacancies in the layered lattice after the first charge. Moreover, the double-layer coated samples exhibited higher rate capacities than the uncoated and single-layer coated samples, attributed to the suppression of undesired SEI layers and the fast charge transfer reaction kinetics. 3.2.4. Inert Li host compounds (ionic conductors) The abovementioned inert coating layers can improve the interphase stability between the electrode and electrolyte. Nevertheless, these inert substances are usually poor electronic and ionic conductors, and the coating layers of these inert metal oxides/fluorides/phosphates often lead to a reduced reversible capacity and a poorer rate discharging ability. In contrast, some Li-contained oxides with Li host structures, for example LiAlO2 [374-377], are more effective for enhancing the electrochemical performance because they have high Li+ conductivity and can provide the tunnels for Li+ transportation during the charge and discharge processes. Markovsky et al. [378] investigated the influence of LiAlO2 coating layers with various thickness on LiNi0.8Co0.15Al0.05O2 electrode by ALD, and found that the 2 nm LiAlO2-coated electrode exhibited ~3 times lower capacity fading and lower potential hysteresis than the uncoated electrode because of the enhanced structure stability. Some other Li+ ion-conducting salts including Li3VO4, Li2ZrO3, Li2MoO4, Li3PO4, Li2SiO3, Li2Si2O5, Li3V2(PO4)3, LiCoPO4, Li3NbO4, Li2TiO3, Li4Ti5O12, Li0.5La0.5TiO3, Li2BO3 and LiMnPO4 are also utilized as coating materials for improving the rate capability of the cathodes [248, 271, 275, 379-391]. Among them, layered Li2TiO3 is a promising coating substance, because its (002) lattice distance of 0.480 nm is similar to the (003) lattice distance of 0.468 nm in the LiNixCoyMnzO2, which can greatly facilitate the Li+ ion transport; Li2TiO3 also has a wide electrochemical window [268]. Zhang’s work [392] proved that Li2TiO3-coated LiNi0.8Co0.15Al0.05O2 particles showed much higher capacity retentions of 94% and 84% than the pristine NCA counterpart (60% after 200 cycles) after 200 and 400 cycles at 0.5 C, respectively. Interestingly, the Li3PO4 coating layers can also scavenge both HF and H2O (usually <50 ppm) in the electrolyte by the reaction formula: Li3PO4+HF (or H2O) → LixHyPO4 (or POxHy)+LiF (or Li2O) [393]. Jo et al. [393] prepared Li3PO4-coated 47

LiNi0.6Co0.2Mn0.2O2 via treatment of the residual Li compounds such as LiOH and Li2CO3 on the surface of LiNi0.6Co0.2Mn0.2O2 with 1 wt% H3PO4 (Fig. 14A-B). The surface-modified electrode showed higher capacity of 161 mAh g–1and capacity retention of 94.1% than the uncoated electrode (127 mAh g–1 and 76.1%) after 150 cycles at 1 C in a potential range of 3.0–4.3 V (Fig. 14C-D), because of the efficient depletion of Li residuals and formation of the high-ion-conductivity Li3PO4 (6×10–8 S cm–1) on the NCM particle surface and the effective inhibition of the HF corrosion and undesired side reactions (Fig. 14E-F).

Fig. 14 (A) Schematic illustration of Li3PO4 coating on LiNi0.6Co0.2Mn0.2O2 particle surface. (B) TEM images, (C) cycling performance at 1 C between 3.0–4.3 V, and (D) rate capability of the bare and coated samples. (E) HF titration and transition metal dissolution results for the electrolyte, and (F) TEM bright field images of the bare and coated NCM (1 wt%) after 150 cycles. Adapted with permission [393]. Copyright 2015 Tsinghua University Press and Springer.

In the similar way, Liu et al. [275] used 1 wt% H2SiO3 to treat LiNi0.6Co0.2Mn0.2O2 for Li2Si2O5-coated LiNi0.6Co0.2Mn0.2O2 particles, and the coated electrode exhibited higher initial capacities of 213.9 and 121.6 mAh g–1 than the uncoated electrode (196.8 and 92.1 mAh g–1) at 0.1 and 10 C, respectively, and also a higher capacity retention of 67% after 200 cycles at 5 C (vs. 52% for the uncoated electrode). Sun et al. [394] investigated the effect of Li2ZrO3 coating on the crystal structure and electrochemical

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property of LiNi0.6Co0.2Mn0.2O2. The Li2ZrO3 coating did not change the layer structure of LiNi0.6Co0.2Mn0.2O2, but the lattice parameters of a and c increased with the Li2ZrO3 content because of the diffusion of Zr4+ ions into the body structure during hightemperature process; 1 wt% Li2ZrO3 coating with 5 nm thickness can effectively increase the Li+ diffusion coefficient and restrain the side reactions with respect to electrolyte, and thus the modified electrode showed high initial capacity of 190 mAh g–1 and capacity retention of 85% after 50 cycles at 0.1 C. Compared with the abovementioned Li host salts, inorganic solid state electrolytes possess obvious advantages of higher Li+ ion conductivity, higher thermal stability, wider electrochemical window, etc. Recently a few inorganic solid electrolytes (e.g., LATP, LLTO, LiPON and LGPS) with high Li+ ion conductivity (10–6–10–2 S cm–1 at room temperature) are chosen as coating substances on cathode particles to improve the electrochemical performance and safety especially at high working temperature [121, 395-401]. Although the introduction of a coating layer creates another interface between the active material and the coating material, which may add an extra resistance to the Li transportation, it is compensated by the fast Li transportation inside the coating layer and the suppression of SEI formation. Therefore, the overall impedance is reduced significantly in the coated sample, and the rate capacity is also enhanced [402]. Choi et al. [403] coated LiNi0.6Mn0.2Co0.2O2 with Li1.3Al0.3Ti1.7(PO4)3 (LATP) by a sol-gel process followed by calcination. The 0.5 wt% LATP-coated electrode showed higher capacity of 164 mAh g–1 and capacity retention of 98% than the pristine electrode (132 mAh g–1 and 87%) after 100 cycles at 1 C in a voltage potential of 3.0–4.3 V, because of the enhancement of Li+ ion movement at the interface of the electrode and protection of the electrode from the electrolyte attack; however, an increase of the coating amount (≥1.0 wt%) retarded the kinetics of the Li+ ion interaction and resulted in a degradation of the cell performance. Liu at al. [404] in-situ coated LiNi0.6Co0.2Mn0.2O2 particles with Li3xLa2/3-xTiO3 (LLTO) by a sol-gel method. The LLTO coating layer can effectively alleviate the structure degradation and suppress the interfacial resistance increase, and thus the 1 wt% LLTO-modified electrode (10 nm thick LLTO layer) exhibited higher capacity of 165 mAh g−1 and retention rate of 91.1% than the pristine electrode (142 mAh g−1 and 78.3%) after 80 cycles at 0.5 C in a voltage range of 3.0−4.5 V. Moreover, the LLTO-coated electrode showed excellent storage stability against moisture and air attack, while the pristine electrode delivered a decreased discharge capacity and poor cycling performance after exposure to air for 2 weeks. The inorganic coating usually require complex and cost-consuming process steps, however, the coating with polymer-based (e.g., PVDF, PAN, PMMA, PPC and PEO) electrolyte on the cathode particles will be more effective on the basis of the polymer electrolyte film-forming capability and the ionic conductivity.[405, 406] Note that a few ionic polymer conductors (e.g., polyimide, PI) are also used as coating material for

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improve the cathode performance [405, 407]. The ionic conductivity can reach 0.15 mS cm–1 by swelling the PI film with liquid electrolyte, which ensures a satisfactory charge/discharge rates [408]. Lee et al. [409] directly coated 20 nm-thick poly(tri(2(acryloyloxy)ethyl) phosphate (PTAEP) gel polymer electrolyte layers on LiNi1/3Co1/3Mn1/3O2 particle surface by a simple UV-assisted polymerization without impairing electronic/ionic conduction pathways performed in the cathode. The ionconductive coating protected the NCM surface from attack of the violent liquid electrolyte and consequently suppressed the harmful side reactions between the NCM and electrolyte. As a result, the coated electrode showed higher capacity of 155 mAh g–1 and capacity retention of 84% than the pristine electrode (155 mAh g–1 and 73%) after 50 cycles at 1 C in a voltage range of 2.8–4.6 V. The coated electrode charged to 4.6 V also yielded less exothermic heat of 649 J g–1 and higher exothermic peak temperature of 294 oC than the pristine electrode (576 J g–1 and 284 oC). 3.2.5. Electronic conductors Similar to the Li host coating materials that can facilitate the Li+ ion transport, some types of electronic conductors that can facilitate the electron transport are also utilized as coating materials for improving the cycling performance of the cathode materials. The coating of electronic conductors can also provide electrons to counteract the cation-insertion-induced charge imbalance upon Li+ ion insertion, allowing for rapid Li+ ion incorporation [410]. Carbon-based materials (particularly amorphous carbon and graphene) are usually preferred because of their low cost, high electrical conductivity and easy processing process [411-425]. Nevertheless, the Li dendrites and SEI films are prone to form on the carbon surface during C-D cycling and thus could result to irreversible capacity loss and even operation safety problem. Hsieh et al. [426] coated ~500 nm LiNi1/3Co1/3Mn1/3O2 particles with 0.08 wt% amorphous carbon with a thickness of ~0.36 nm by carbonization of glucose during an in-situ sintering process. The carbon coating can not only reduce the dissolution of transition metals but also assist the ionic diffusion in the solid solution and electron jumping across the electrode, and thus the carbon-coated electrode showed higher capacity of 130 mAh g–1 and retention rate of 94.2% than the pristine electrode (65 mAh g–1 and 51.7%) after 50 cycles at 1 C in a voltage range of 2.8–4.5 V. Zhang et al. [427] coated LiNi0.8Co0.15Al0.05O2 particles with carbon nanotubes (CNTs) by simple mechanical grinding process without damaging the crystal structure and morphology. The CNT-modified NCA electrode exhibited higher capacity of 181 mAh g–1 and retention rate of 96% than the bare NCA electrode (153 mAh g–1 and 90%) after 60 cycles at 0.25 C, because the CNT coating cannot only increase the electrical conductivity but also inhibit the side reactions with the electrolyte. Shim et al. [428] first reported the reduced graphene oxide-encapsulated LiNi0.6Co0.2Mn0.2O2 secondary particles by a wet chemical method with the assistance of (3-

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aminopropyl)triethoxysilane (APTES) (Fig. 15A-D). The ~2 nm thick graphene layers protected the particles from direct contact with the electrolyte, inhibited the oxygen release and detrimental side reactions, and suppressed the structural transition from the rhombohedral layered structure to the spinel or even rock-salt phase (Fig. 15F). The graphene coating also significantly facilitated the electron/Li+ transfer and reduced the charge/discharge polarization during cycling. Consequently, the graphene-modified electrode showed a higher capacity of ~160 mAh g–1 than the pristine electrode (~110 mAh g–1) after 100 cycles at 1 C in a voltage range of 3.0–4.5 V (Fig. 15E), and a higher rate capacity of 133 mAh g–1 than the pristine electrode (105 mAh g–1) at 10 C. The graphene-coated sample also exhibited an higher exothermic peak temperature of 219.2 °C than the bare sample (285.2 °C), which indicated that the thermal decomposition was significantly delayed by means of the graphene wrapping (Fig. 15G).

Fig. 15 (A) Schematic illustration of a reduced graphene oxide-coated NCM particle (rGO-NCM), (B-C) SEM and (D) TEM images of the rGO-NCM particles, and (E) voltage-capacity curves of both the NCM samples at 1 C, (F) HAADF STEM images of (a) NCM and (b) rGO-NCM after 100 cycles, and (G) DSC profiles of the NCM and rGO-NCM charged to 4.5 V. Adapted with permission [428]. Copyright 2017 American Chemical Society.

The electrochemical property of Ag-coated LiNi1/3Co1/3Mn1/3O2 has also been investigated though the high cost of Ag coating [429]. The Ag coating can increase the conductivity and lower the polarization of the cell, and promote the formation of compact and protective SEI layer. Therefore, the 6.6 wt% Ag-coated electrode showed higher capacity of 160 mAh g–1 and retention rate of 94.7% than the pristine electrode (143 mAh g–1 and 86.7%) after 50 cycles at 20 mA g–1 in a voltage range of 2.8–4.4 V. However, the Ag coating is unstable in the non-aqueous electrolyte system during the charge/discharge processes, because it can dissolve into the electrolyte or be oxidized to silver oxide [430].

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3.2.6. Dual/hybrid conductors The surface modification of cathode materials with conductive polymers such as polypyrrole (PPy), polyaniline (PANi), polyethylene glycol (PEG) and poly(3,4ethylenedioxythiophene) (PEDOT) is quite beneficial to improving the cycling performance without sacrificing the original capacity, owing to their high electronic conductivity and electrochemical stability [431, 432]. However, they are not ionic conductive. So it is highly desirable to apply dual-conductive polymers with both high Li+ ion and electron conductivities (e.g., PEDOT-PSS and PEDOT-PEG), aiming to greatly enhance both the Li+/e– charge transfer on the electrode surface [432-434]. Ju et al. [432] coated LiNi0.6Co0.2Mn0.2O2 particles with a dual-conductive PEDOT-PEG copolymer by a wet-chemical method. The 11–25 nm thick films on the electrode surfaces made the surface-modified NCM particles exhibit higher electronic conductivity of 0.2 S cm–1 than the pristine particles (1.6×10–6 S cm–1). Additionally, the ionic conductivity of the PEDOT-PEG film soaked with the liquid electrolyte (1 M LiPF6 in EC/DMC) was 4.2×10–3 S cm–1, indicating fast ion transport through the thin surface layer. Besides, the conducting polymer layer formed on the cathode can suppress the growth of a resistive layer and inhibit the dissolution of transition metals from the active cathode materials. Consequently, the PEDOT-PEG-coated electrode showed a higher capacity of 172 mAh g–1 than the uncoated (168 mAh g–1) and PEDOT-coated (160 mAh g–1) electrodes after 100 cycles at 0.5 C. Especially after 100 cycles at 55 oC, the PEDOT-PEG-coated electrode showed a higher capacity of 167 mAh g–1 than the pristine electrode (90 mAh g–1). Coating the cathode particles with two different materials as hybrid shell layers has also been developed. Makhonina et al. [298] coated LiNi0.40Mn0.40Co0.20O2 particles with both amorphous carbon and Al2O3 with the C-Al2O3 coating layer of 15–25 nm in thickness. The C-Al2O3-coated electrode showed higher capacity of 91 mAh g–1 and retention rate of 90.9% than the Al2O3-coated electrode (83 mAh g–1 and 83.1%) after 110 cycles at 10 Ah g–1 in a potential range of 3.0–4.5 V. The C-Al2O3-coated electrode also had a superior rate performance over the pristine and Al2O3-coated electrodes. The improvement in the cycling performance were ascribed to the following two reasons: (1) the Al2O3 coating can improve the surface stability of the cathode materials by suppressing the detrimental chemical side reactions between the electrode surface and electrolyte by acting as the physical barrier or/and as a scavenger for HF and water; (2) the amorphous carbon film should increase the conductivity of the mixed coating layer and prevent the loss of the electrical contact. Kim et al. [435] coated reduced graphene oxide (rGO) and AlPO4 on the surfaces of Li1.190Mn0.540Co0.143Ni0.127O2 particles by an electrostatic interaction between the coating precursors and the active electrode particles with a control of pH. The ~4 nm thick rGO-AlPO4 hybrid film on the cathode particle surface can decrease the Li content in the SEI layer and facilitate the fast charge transfer,

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and thus the 1wt%rGO-2wt%AlPO4-coated electrode showed a higher capacity of 205 mAh g–1 than the pristine (190 mAh g–1 after 100 cycles) and 2 wt%AlPO4-coated (190 mAh g–1) electrodes after 150 cycles at 0.1 C in a voltage range of 3.0–4.8 V. Even after 150 cycles at 55 oC, the 1wt%rGO-2wt%AlPO4-coated electrode showed a higher capacity of 233 mAh g–1 than the pristine (133 mAh g–1 after 100 cycles) and 2wt%AlPO4-coated (215 mAh g–1) electrodes. Coating cathode materials with hybrid layers containing both electronic and ionic conductors would be a promising technology in the future, based on the compensable advantages of the electronic and ionic conductors [436-438]. Zhou et al. [438] synthesized 200 nm Li3V2(PO4)3 particles and then co-coated them with amorphous carbon and Li7La3Zr2O12 (C-LLZO/LVP) through a sol-gel method. The carbon coating acted as an electron conductor to enhance electron transport, while the Li7La3Zr2O12 acted as a fast ionic conductor to facilitate Li+ ion transport on the surface of Li3V2(PO4)3. Therefore, the C-LLZO/LVP electrode showed higher capacity of 176 mAh g–1 and retention rate of 99.3% than the C/LVP electrode (169 mAh g–1 and 97.7%) after 20 cycles at 1 C between 3.0–4.8 V. Even after 300 cycles at 10 C, the CLLZO/LVP electrode showed higher capacity of 139 mAh g–1 and retention rate of 92.6% than the C/LVP electrode (103 mAh g–1 and 82.9%). Dang et al. [439] coated LiNi1/3Co1/3Mn1/3O2 particles with both amorphous carbon and Li2SiO3 through sol-gel and carbonization. The Li2SiO3-C dual functional coating layer with 4–8 nm thickness can protect the cathode material from the electrolyte corrosion and avoid unfavorable the interfacial side reactions. Besides, Li2SiO3 and carbon acted as Li-ion conductor and electron conductor respectively, and promoted Li+ ion and electron migration. Therefore, the Li2SiO3-C-coated electrode showed higher capacity of 130 mAh g–1 and retention rate of 95.0% than the pristine (118 mAh g–1 and 92.0%) and Li2SiO3-coated (121 mAh g–1 and 93.7%) electrodes after 90 cycles at 1 C in a voltage range of 3.0–4.3 V. Even after 500 cycles at 5 C, the Li2SiO3-C-coated electrode showed higher capacity of 96 mAh g–1 and retention rate of 91.4% than the pristine (72 mAh g–1 and 70.6% after 450 cycles) and Li2SiO3-coated (82 mAh g–1 and 79.6%) electrodes. Unlike the single shelled structure, recently another special core-shell heterostructure with double shell layers on the cathode core have also been exploited. Xia et al. [440] fabricated a layered@spinel@carbon cathode particle by in situ synchronous carbonization-reduction of a polydopamine-coated layered Li1.14Ni0.19Co0.19Mn0.47O2 particle. During the carbonization-reduction process, a small portion of the layered structure was gradually transformed into the spinel structure, forming the heterostructure with a Li-rich layered oxide core, a spinel interlayer and a carbon outermost layer (Fig. 16A-D). This structure combines the advantages of the Lirich layered core and the spinel and carbon shells: (1) the core region offers high capacity; (2) the spinel interlayer possesses 3D Li+ ion diffusion channels to facilitate

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rapid Li+ ion transport and improve the rate capacity; (3) the outermost conductive carbon layer acts as a protection layer and can reduce the HF attack, suppress the transition metal ion dissolution, hinder the side reactions and improve the structural stability and electronic conductivity. As a consequence, the heterostructured electrode maintained a higher capacity of 215 mAh g–1 than the pristine electrode (155 mAh g–1) after 50 cycles at 1 C in a voltage range of 2.0–4.8 V. As the charge-discharge rate increased from 2 C to 20 C, the heterostructured sample also exhibited superior rate performance compared with the pristine sample (Fig. 16E).

Fig. 16 (A) Schematic illustration of the coating mechanism for the protection of the NCM particle with layered@spinel@carbon configuration (L@S@C), and (B-C) TEM images, (D) XRD patterns and (E) cycling performance of the L@S@C samples at different discharge rates between 2.0–4.8 V. Adapted from permission [440]. Copyright 2015 American Chemical Society.

3.2.7. Cathode materials (including concentration gradient) Although the inert coatings can improve the electrochemical and thermal performance, the inert materials do not contribute the capacity. Moreover, such coating layers do not seem to suppress the structural change during the long-term electrochemical cycling, because of the disadvantageous characteristics of the thin layers of only several (or tens of) nanometers in thickness and the amorphous/lowcrystalline phase [441]. Therefore, a few electrochemically-active electrode materials such as LiCoO2, LiMn2O4, LiMnPO4, LiFePO4 and LiCoPO4 are utilized as coating substances to form core-shell structures for increasing the overall electrode capacity [284, 442-450]. In addition to the single shell-based configuration, the core-shell structures with double shells are also adopted [451, 452]. As for the typical core-shell

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structures, there are obvious interfaces between the inner cores and the outer shells. Both the Ni content in the inner cores and Mn content in the outer shells are usually much higher, which is conducive to improving both the cycling performance and thermal stability [152, 453-456]. Liu et al. [450, 457] also found that the LiCoO2-coated LiNi0.8Co0.15Al0.05O2 exhibited better storage stability than the pristine LiNi0.8Co0.15Al0.05O2 due to the better resistance against H2O and CO2. It should be noted that the coating with other commercial cathode materials will also inherit their drawbacks (e.g., the coating of manganese spinel may reduce the calendar life due to the loose of Mn ions during cycling). Sun et al. [441] firstly reported the core-shell structured LiNi0.8Co0.1Mn0.1O2LiNi0.5Mn0.5O2 (or Li[(Ni0.8Co0.1Mn0.1)0.8(Ni0.5Mn0.5)0.2]O2) microparticles with excellent cycling stability and safety in 2005 (Fig. 17A-D). The microparticles of ~15 μm in diameter with 1–1.5 μm thick shells can combine the high capacity derived from the Ni-rich LiNi0.8Co0.1Mn0.1O2 cores and the high thermal stability stemmed from the LiNi0.5Mn0.5O2 shells, and thus the pouch-type LiNi0.8Co0.1Mn0.1O2LiNi0.5Mn0.5O2/graphite full-cell showed higher capacity of 113 mAh g–1 and retention rate of 98% than the LiNi0.8Co0.1Mn0.1O2/graphite battery (94 mAh g–1 and 81%) after 500 cycles at 1 C in a voltage range of 3.0–4.3 V (Fig. 17E). Furthermore, the LiNi0.8Co0.1Mn0.1O2-LiNi0.5Mn0.5O2 particles charged to 4.3 V exhibited less exothermic heat of 2261 J g–1 at ~250 oC than the LiNi0.8Co0.1Mn0.1O2 particles (3285 J g–1 at ~220 oC), indicating the significant improvement in thermal stability owing to the effective suppression of the oxygen evolution from the highly-delithiated Li1–xNi0.8Co0.1Mn0.1O2 by the thermally-stable outer LiNi0.5Mn0.5O2 shells (Fig. 17F).

Fig. 17 (A) Schematic illustration of the preparation process of core-shell structured Li[(Ni0.8Co0.1Mn0.1)1–x(Ni0.5Mn0.5)x]O2 microparticles, (B-D) SEM and EDS images of the coreshelled Li[(Ni0.8Co0.1Mn0.1)0.8(Ni0.5Mn0.5)0.2]O2 particles, and (E) cycling performance at 1 C between 3.0–4.3 V and (F) DSC curves of the bare and core-shelled LiNi0.8Co0.1Mn0.1O2 samples charged at 4.3 V. Adapted with permission [441]. Copyright 2005 American Chemical Society.

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Inspired by the Sun’s works, different core-shell structures are fabricated [116, 457]. Du et al. [458] coated layered LiNi0.80Co0.15Al0.05O2 (a=2.8602 Å, c=14.2239 Å) with layered LiNi1/3Co1/3Mn1/3O2 with similar lattice parameters (a=2.8633 Å, c=14.165 Å) to achieve their structural compatibility, and thereby the modified electrode showed much higher capacity retentions of 96.2% and 84.3% after 100 cycles at 0.2 C at 25 and 50 oC than the uncoated electrode (86.9% after 100 cycles at 25 oC and 68.8% after 50 cycles at 50 oC), respectively. Lee et al. [459] coated LiNi0.80Co0.15Al0.05O2 with LiFePO4 NPs (~500 nm) to improve the cycling performance by a co-precipitation method followed by ball-milling and calcination. The coated electrode exhibited a high initial discharge capacity of 163 mAh g–1 without capacity dropping at 25 oC, and presented a remarkably higher cycling retention of 89% after 400 cycles at 50 oC than the uncoated electrode (82% after 250 cycles). Noh et al. [460] prepared Li2MnO3coated LiNi0.5Co0.2Mn0.3O2 particles via ball-milling and subsequent sintering, and the coated electrode exhibited excellent cycling stability (215 mAh g–1 after 20 cycles at 0.1 C) owing to the improved structural and thermal stability by the Li2MnO3 layers. Though the great improvement in cycling stability by applying the core-shell configuration, there is a structural mismatch between the core and the shell [461-464]. For example, the Ni-rich NCM core usually experiences a larger volume variation of 9– 10% during Li insertion/extraction, whereas the LiNi0.5Mn0.5O2 shell has a smaller volume change of 2–3% in lattice cell volume. Owing to the different volume change (or expansion/shrinkage) between the core and the shell during lithiation/delithiation, large voids of tens of nanometers at the core/shell interface would generate in the coreshelled particles upon cycling due to the inner mechanical stress/strain, which could result in gradual separation between the core and the shell, inhibit the Li+ ion and electron transfer, and eventually lead to a drastic decline of the battery performance. To overcome this issue, a new design concept was exploited by the Cho group [465] via nanoscale surface treatment of the cathode particles, exactly for the primary particle surfaces (Fig. 18A). The resultant primary particle surfaces had a higher cobalt content and a cation-mixing phase (Fm3̅m) with nanoscale thickness (~10 nm, as a surface pillar layer) in the LiNi0.6Co0.2Mn0.2O2 particles (Fig. 18C-F), which can effectively mitigate the micro-cracks in the electrode by suppressing the structural change from the layered to the rock-salt phase during the electrochemical cycling. As a result, the modified electrode showed a higher capacity retention of 80% (from 195 to 156 mAh g–1) than the pristine electrode (65% retention from 198 to 130 mAh g–1) after 150 cycles at 1 C in a voltage range of 3.0–4.45 V at 60 oC (Fig. 18B). SEM and STEM characterizations of the cycled electrodes further confirmed that the modified sample had less microcracks and mixed spinel-like and rock-salt phases than the pristine sample (Fig. 18G-J). DSC measurements of the charged electrodes at 4.45 V showed that the surface modified sample showed higher main peak temperature of 264 oC and less heat

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generation of 696.8 J g–1 of than the pristine sample (230 oC and 910.9 J g–1), suggesting that the nanoscale surface treatment can greatly enhance the thermal stability and safety of the active cathode materials.

Fig. 18 (A) Schematic illustration of the surface treatment of the primary NCM particles (ST-NCM) for cation-mixing phase on the particle surfaces, (B) cycling performance of the pristine and surfacetreated NCM samples at 1 C between 3.0–4.45 V at 60 oC, (C) SEM, (D) STEM and (E-F) enlarged STEM images in (D) of the ST-NCM particles, and cross-sectional FIB SEM and STEM images of the (G-H) pristine and (I-J) surface-treated NCM particles after 150 cycles. The insets of (H) and (J) show the corresponding FFT patterns. Adapted with permission [465]. Copyright 2015 American Chemical Society.

Researchers have also combined concentration-gradient approaches in the shell or core with the core-shell design to ensure a smooth transition in composition from the inner core to the outer shell surface to solve the problems associated with the core-shell mismatch in the core-shell configuration [53, 462-464, 466, 467]. A core-concentration gradient-shelled LiNi0.95Co0.026Mn0.024O2 (average composition, ~11 μm in size) particle was prepared by Jun et al. [468] via heating a core-shelled precursor. The hightemperature lithiation of the precursor led to a continuous and smooth concentration gradient from the bulk to the shell due to the inter-diffusion of the cations, and the resulted LiNi0.95Co0.026Mn0.024O2 particle with a ~8 μm LiNiO2 core, a ~1 μm-thick transition shell layer and a ~0.5 μm-thick outer shell of LiNi0.87Co0.065Mn0.065O2 (Fig.

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19A-C). The core-concentration gradient-shell structure could inhibit the phase _

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transition from R3m (layered structure) to Fm3m (NiO-like cubic phase) caused by the cation mixing upon cycling, and maintain the structural stability of the high-Ni core with the Ni-depleted protective shell layers (Fig. 19F-H). Furthermore, the concentration gradient within the shell would effectively relax the core/shell interface strain that could have formed due to the sharp composition change for improving the long-term cycling of the cathode. Therefore, the LiNi0.95Co0.026Mn0.024O2/Li half-cell showed higher capacity of 199 mAh g–1 and retention rate of 90% than the LiNiO2/Li half-cell (161 mAh g–1 and 74%) after 100 cycles at 0.5 C in a potential range of 2.7– 4.3 V (Fig. 19D). Additionally, the pouch-type LiNi0.95Co0.026Mn0.024O2/graphite fullcell showed higher capacity of 163 mAh g–1 and retention rate of 84% than the LiNiO2/graphite full-cell (101 mAh g–1 and 50%) after 1000 cycles at 1 C in a potential range of 3.0–4.2 V (Fig. 19E). Guo et al. [469] reported core-concentration gradientshelled particles with various element composition from the outer shell of LiNi1/3Co1/3Mn1/3O2 to the inner core of LiNi0.8Co0.15Al0.05O2. The cathode particle combined the merits of the concentration-gradient protecting buffer of NCM and the inner core of NCA, and thus exhibited higher capacity retention (99.8% after 200 cycles at 0.5 C) and improved thermal and air stability (onset temperature of 250.5 oC and heat of 1237 J g–1) than the pristine NCA (59.3%, 191.4 oC and 1472 J g–1).

Fig. 19 (A-B) SEM images and (C) EPMA compositional linear scan of core-shelled NCM particles, cycling performance of the (D) half-cells at 0.5 C between 2.7–4.3 V and (E) full-cells with graphite anodes at 1 C between 3.0–4.2 V, and TEM images, corresponding electron diffraction spectra (EDS) and FFT patterns of the (F-G) core-shelled NCM and (H) LiNiO2 after 100 cycles. Adapted with permission [468]. Copyright 2017 American Chemical Society.

To simplify the preparation process of the core-shelled precursors and cathode particles, Zhang et al. [470] proposed a novel strategy to in situ generate an integrated

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hierarchical spinel layer on the layered LiNi0.8Co0.1Mn0.1O2 (SC-LNCMO) particle surface through a pH modulation-induced gradient change of Mn ion valences in the precursor. The self-forming outer spinel shell layer was tightly integrated onto the inner layered phase core by atom-scale interfacial junctions of a Ni concentration-gradient transition phase shell layer, which can effectively reduce the core-shell phase mismatch. The core-concentration gradient-shell structure can effectively improve the structure stability and stabilize the formation of stable SEI films during long-term cycling. As a result, the SC-LNCMO electrode exhibited higher capacity of 159 mAh g–1 and retention rate of 93% than the LNCMO electrode (91 mAh g–1 and 63%) after 100 cycles at 1 C in a potential range of 3.0–4.5 V. The SC-LNCMO electrode also showed higher capacity of 146 mAh g–1 and retention rate of 93% than the LNCMO electrode (40 mAh g–1 and 43%) at –20 oC after 100 cycles at 0.1 C. Even after a storage of 90 days, the SC-LNCMO electrode displayed a slighter Li+/Ni2+ mixing increase from 1.94% to 2.44% (vs. 3.26% to 7.02% for LNCMO) and better cycling stability. Another concentration gradient core-shelled LiNi0.5Co0.2Mn0.3O2 (CG-NCM, average composition) particle was fabricated by Song et al. [463] through calcination of a double-shelled [(Ni0.8Co0.1Mn0.1)2/7(Ni1/3Co1/3Mn1/3)3/14(Ni0.4Co0.2Mn0.4)1/2](OH)2 precursor. The CG-NCM particle consisted of a ~10 μm core with the gradient change of Ni and Mn concentrations but constant Co concentration toward the shell, a ~1 μmthick transition phase shell, and a ~1 μm-thick outer shell. Because of the special particle configuration, the CG-NCM electrode showed a higher capacity of 188.8 mAh g–1 than the normal NCM electrode (175.2 mAh g–1) at 55 oC after 100 cycles at 0.5 C in a potential range of 3.0–4.5 V. The CG-NCM particle also generated less heat of 609.9 J g–1 at 302.1 oC than the normal NCM particle (972.6 J g–1 at 290.2 oC). Despite the partial alleviation of the structure and expansion/shrinkage mismatches by the combination of the concentration gradient design in the core or shell, the low Mn or Al content in the thin outer shell (usually ≤2 μm) cannot effectively stabilize the particle surface especially during high-temperature cycling [471]. Even provided that the shell is further thickened, the thermal stability and cycling stability can be improved, but the total energy density would decrease [40]. To further resolve the problems caused by the sharp composition difference in the core and the shell for enhancing the cycling performance, researchers have also developed a full-concentration-gradient (FCG) structure to obtain high-performance Ni-rich cathode particles, where the Ni concentration decreases gradually whereas the Mn or Al concentration increases gradually from the center to the outer layer of each particle [471-476]. The FCG design assures a continuous and smooth transition from the surface to the bulk and enables the FCG cathode particle to take advantage of high capacity delivered by the Ni-enriched core and simultaneously the thermal and cycling stability ensured by the Mn- or Alenriched surface layer [473, 476-479].

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The Sun group [471] prepared FCG LiNi0.75Co0.10Mn0.15O2 (average composition) particles via co-precipitation and calcination involving the precipitation of transitionmetal hydroxides from the precursor solutions with the continuous concentration change of Ni/Co/Mn with the reaction time. The FCG particles had smooth Ni/Mn concentration change but constant Co concentration toward the particle surface, and combined the advantages of the high capacity from the higher Ni content in the center and the high thermal stability from the higher Mn content on the surface (Fig. 20A-D). Therefore, the coin-type FCG-LiNi0.75Co0.10Mn0.15O2/Li half-cell showed a higher capacity of 190 mAh g–1 than the conventional LiNi0.86Co0.10Mn0.04O2 (inner composition of the FCG particles) and LiNi0.70Co0.10Mn0.20O2 (outer composition) halfcells (118 and 180 mAh g–1) after 100 cycles at 0.2 C in a potential range of 2.7–4.5 V (Fig. 20E). Moreover, the pouch-type FCG-LiNi0.75Co0.10Mn0.15O2/graphite full-cell exhibited high capacities of 29 and 24 mAh with capacity retentions of 91% and 69% after 1000 cycles at 1 C in a potential range of 3.0–4.2 V at 25 and 55 oC, respectively (Fig. 20F). Zhang et al. [480] synthesized LiNi0.8Co0.15Al0.05O2 particles with decreasing Ni and Co concentration and increasing Al concentration toward the particle surface. Due to the effective protection of the inner components from direct attacking by the electrolyte and suppression of the side reactions on the particle surface between Ni4+ ions and electrolyte, the FCG NCA electrode showed higher capacity retention of 93.6% after 100 cycles and thermal stability than the pristine NCA electrode.

Fig. 20 (A) Schematic diagram of a full-concentration-gradient (FCG) NCM particle, (B) SEM element mapping and (C) EPMA line scan of the integrated atomic ratio of transition metals as a function of the distance from the particle center to the surface, (D) TEM image of the local structure near the particle edge showing a highly-aligned nanorod network, and cycling performance of the (E) NCM/Li half-cells and (F) NCM/graphite full-cells. Adapted with permission [471]. copyright 2012 Macmillan.

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Different from the constant concentration changes of Ni/Mn in the FCG particles, the Sun group [478] further designed and fabricated two-sloped-full-concentrationgradient (TSFCG) LiNi0.85Co0.05Mn0.10O2 particles by a stepwise co-precipitation method to achieve maximum possible discharge capacity by having a Ni-enriched core and to simultaneously ensure high chemical and thermal stability by having an outer Mn-enriched layer. The coin-type TSFCG/Li half-cell showed higher capacity of ~190 mAh g–1 and capacity retention of 92% than the conventional LiNi0.85Co0.11Al0.04O2 (150 mAh g–1 and 75%) after 100 cycles at 0.5 C in a potential range of 2.7–4.3 V. Even at higher temperature of 55 oC or higher charge-discharge rates, the TSFCG cathode showed excellent cycle performance. Furthermore, the TSFCG/CNT-Si full-cell delivered higher capacity of ~200 mAh g–1 and capacity retention of 81% with an average Coulombic efficiency of 99.8% after 500 cycles at 1 C in a voltage range of 2.0–4.15 V. The energy density was calculated to be ~726 Wh kg–1 (average voltage of 3.63 V), which translated to a high energy density of 327 Wh kg–1 for the TSFCG/CNTSi full-cell based on the 45 wt% mass fraction of the cathode materials in a commercial 18650 cell, satisfying the energy density limit by the driving range requirement (~300 miles) for electric vehicles. Their recent studies also disclosed that the preheating and sintering conditions were critical in controlling the phase separation in the final product, getting rid of the detrimental Li2CO3 for a longer life, and reducing the occupancy of Ni in the Li layer for a better cycling performance [481]. Unlike the abovementioned FCG structure with the gradient change of Ni and Mn without the Co gradient change from the particle center to the surface, another FCG structure with the Ni and Co gradient change but constant Mn concentration has also been developed. Ju et al. [482] prepared LiNixCo0.16Mn0.84−xO2 particles (x=0.64, 0.59 or 0.51, average composition) with varying concentration gradients of Ni and Co from the particle center (0.62–0.74 mol% Ni and 0.05 mol% Co) to the surface (0.48–0.62 mol% Ni and 0.18 mol% Co), i.e., FCG with fixed Mn concentrations (20, 25 or 33 mol%). An increase in Ni content improved the capacity, whereas a higher amount of Mn delivered better capacity retention and thermal properties at the expense of capacity and rate performance. Therefore, the electrode with an optimized composition of LiNi0.59Co0.16Mn0.25O2 showed the highest capacity of 188 mAh g–1 with capacity retention of 96% after 100 cycles at 0.5 C in a potential range of 2.7–4.3 V. They also fabricated FCG LiNi0.60Co0.15Mn0.25O2 particles (average composition) with fixed Mn concentration but gradient changed Ni and Co concentrations by co-precipitation [483]. The constant Mn concentration across the particle provided outstanding cycle life and safety, the linear decrease in Ni concentration toward the particle outer surface resulted in high capacity, and the gradual increase of the Co concentration from the particle center to the surface reduced the cation mixing and improved the electronic conductivity

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and rate capability. As a consequence, the coin-type FCG-LiNi0.60Co0.15Mn0.25O2/Li half-cell showed higher capacity of 185 mAh g–1 and capacity retention of 91% than the conventional half-cell (163 mAh g–1 and 83%) after 100 cycles at 0.5 C in a potential range of 2.7–4.5 V at 55 oC. The pouch-type FCG-LiNi0.60Co0.15Mn0.25O2/graphite fullcell also exhibited higher capacities of 166 and 135 mAh g–1 with capacity retentions of 88% and 70% after 1000 cycles at 1 C in a potential range of 3.0–4.4 V at 25 and 55 oC, respectively. The full-concentration-gradient method has been further developed with the composition gradient change of Mn and Co but constant concentration of Ni, in order to greatly increase the capacity and remain the electrochemical and thermal stabilities [484]. Noh et al. [485] prepared FCG LiNi0.80Co0.06Mn0.14O2 (average composition) particles with a continuous compositional change from LiNi0.80Co0.20O2 (center) to LiNi0.80Co0.01Mn0.19O2 (surface) by co-precipitation and following sintering. The FCGLiNi0.80Co0.06Mn0.14O2 electrode showed higher capacity retentions of 87% and 81.4% than the conventional bulk LiNi0.80Co0.06Mn0.14O2 electrode after 100 cycles at 0.5 C in a voltage range of 2.7–4.3 V at 25 and 55 oC, respectively. Moreover, the FCGLiNi0.80Co0.06Mn0.14O2 particles charged to 4.3 V exhibited higher main exothermic temperature of 250 oC and lower heat of 810 J g–1 than the conventional LiNi0.80Co0.20O2 particles (230 oC and 984 J g–1), indicating that the presence of Mn4+ in the surface delayed the phase transformation and resulted in less evolution of oxygen from the structure due to the stronger Mn4+-O bond in the surface. It should be mentioned that the synthesis of the FCG particles usually including the multistep co-precipitation processes requires additional tanks for offering transition metals and may increase the total cost and time. To tackle this issue, Liao et al. [486] prepared FCG LiNi0.76Co0.10Mn0.14O2 particles by lithiating core/double-shelled [Ni0.9Co0.1]0.4[Ni0.7Co0.1Mn0.2]0.5[Ni0.5Co0.1Mn0.4]0.1(OH)2 precursors at 750 °C for 20 h in a flowing oxygen atmosphere. In the FCG particles, the Ni concentration decreased linearly, the Mn concentration increased gradually, and the Co concentration remained constant from the center to the surface of each particle, which could make the particles keep excellent thermal stability, high energy density and good cycling stability. Specifically, the coin-type LiNi0.76Co0.10Mn0.14O2/Li half-cell battery (active material: 15–20 mg cm–2, areal capacity: 3–4 mAh cm–2) showed high capacities of 197 and 181 mAh g–1 and retention rates of 99% and 100% after 100 cycles at C/3 and 1 C, respectively. Moreover, the LiNi0.76Co0.10Mn0.14O2/graphite full-cell showed high capacity of 19.6 mAh and retention rate of 89% after 500 cycles at C/3. Another simplified route for FCG cathode particles was very recently conceived by the Cho group [487]. They fabricated FCG LiNi0.6Co0.2Mn0.2O2 particles by a polystyrene bead (PSB) cluster-incorporated co-precipitation method without the additional tanks (Fig. 21A-B). During the following annealing process, the decomposition of the PSBs in the

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precursor cores led to the formation of voids, which would enhance the electrolyte wettability and buffer the volume variation upon cycling. Accordingly, the FCG samples had a slightingly lower tap density of ~2.4 g cm–3 than the pristine sample (~2.6 g cm–3). Simultaneously, the Ni3+ ions in the precursors were reduced to stable Ni2+ ions on the surface of the active materials during the calcination, which induced the diffusion of other metal ions from the surface to the core for balancing the charge neutrality, resulting in the concentration gradients in the particles for improving the structure and cycling stabilities (Fig. 21C). This unique structure allowed the cycling stability at a high voltage of 4.45 V without any electrolyte additive. The FCG electrode with an electrode loading density of ~2.7 g cm–3 kept ~5% higher discharge capacity (157 vs. 143 mAh g–1) than the pristine electrode after 250 cycles at 1 C in a voltage range of 3.0–4.45 V at 25 oC (Fig. 21D). Furthermore, the FCG electrode exhibited higher capacity of 175 mAh g–1 and capacity retention of 86% than the pristine electrode (133 mAh g–1 and 67%) at 1 C after 250 cycles at 60 °C (Fig. 21E). Compared to the pristine sample after the long-term cycling, the FCG sample maintained the initial porous morphology without micro-crack evolution and the robust layered structure without severe lattice deterioration to NiO rock-salt and cation-mixed phases (Fig. 21F-G).

Fig. 21 (A) Schematic drawing of the synthesis of FCG NCM particles (denoted as PSB-NCM) using PSB and CTAB as sacrificial templates and surfactants, respectively, (B) SEM image of the

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PSB-NCM particle, (C) concentration of the transition metals toward the PSB-NCM particle center with the illustration of the difference of the surface gradient of Ni oxidation state and transition metal concentration, cycling performance of the pristine and FCG NCM samples at 0.5 C charge rate and 1 C discharge rate between 3.0–4.45 V at (D) 25 and (E) 60 oC, respectively, and HR-TEM images and corresponding FFT patterns of the (F) pristine and (G) FCG NCM samples after 250 cycles at 60 oC. Adapted with permission [487]. Copyright 2017 Wiley-VCH.

3.2.8. Compositing with multifunctional materials Apart from the aforesaid body-doping and surface-coating methods, fabricating composites with the cathode particles and multi-functional materials offers another important avenue for improving the battery performance by increasing the structure stability and conductivity of the NCM/NCA-based cathodes [417, 488, 489]. The multifunctional matrices posses high mechanical strength and electrical conductivity, as well as high electrochemical stability during lithiation/delithiation reactions, and can function as mechanical and electrical supports to alleviate the internal stress resulted from the morphological change of the electrode materials upon cycling and facilitate the transport of charges [162, 410, 490, 491]. For example, nitric acid-treated Super P carbon black powders with the surfaces changed from hydrophobic to hydrophilic were dispersed in cathode nanoparticle aqueous solution, sprayed and dried at 300 oC in inert atmosphere for high-performance composite cathode particles [492]. Carbon-based nanomaterials such as carbon nanotubes (CNTs) and graphene with large specific surface area, high electrical conductivity and excellent mechanical flexibility are usually composited with electrode materials for enhancing the cycling performance of the composite electrodes [417, 428, 489, 493-498]. In addition to providing paths for electrical conduction, nano-scale fibers can also be used to disperse the cathode particles for increased contact with the electrolyte [499]. Yoon et al. [500] prepared a LiNi0.8Co0.15Al0.05O2/graphene composite by a high energy ball-milling process. Due to the improved electrical conductivity by the graphene matrix, the composite electrode showed higher capacity of 180 mAh g–1 and retention rate of 97% than the pristine NCA electrode (172 mAh g–1 and 91%) after 80 cycles at 0.3 C within 3.0–4.3 V. The composite electrode also displayed high rate capacities of 152 and 112 mAh g–1 at 10 and 20 C, respectively. Jan et al. [501] prepared a LiNi0.8Co0.1Mn0.1O2/graphene composite by a facile chemical approach (Fig. 22A-D). The 100–200 nm NCM particles were uniformly dispersed on the graphene surface and meanwhile enwrapped by the graphene nanosheet to from a three-dimensional network, which can suppress the nanoparticle agglomeration, avoid the direct exposure of the nanoparticle to the electrolyte to prohibit the side reactions and the active material dissolution, and offer fast electron transfer to improve the electrical conductivity of the electrode. As a result, the composite electrode showed higher capacity of 168 mAh g–1 and retention rate of 92.2% than the pristine LiNi0.8Co0.1Mn0.1O2 electrode (133 mAh g–

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and 76.5%) after 150 cycles at 1 C in a voltage range of 2.7–4.3 V (Fig. 22E-G). The composite electrode also exhibited higher discharge capacity of 164 mAh g–1 and Coulombic efficiency of 95.5% than the pristine electrode (133 mAh g–1 and 90.1%) at 5 C.

Fig. 22 (A) Schematic illustration for preparing NCM/graphene (denoted as G-NCM) composite, (B) SEM and (C-D) TEM images of the NCM/graphene composite, and (E) rate capability, (F) EIS spectra after 5 cycles at 0.1 C and (G) cycling performance at 1 C between 2.7–4.3 V of the bare NCM and composite samples. Adapted with permission [501]. Copyright 2014 Elsevier.

Besides the compositing with electronic conductors, the compositing with ionic conductors (i.e., solid state electrolytes) is also attracting much attention. Nakamura et al. [395] prepared a LiCoO2/Li1.3Al0.3Ti1.7(PO4)3 (LATP) composite granule as low-cost cathode by simple mechanical mill method for all-solid-state Li-ion batteries. The mechanical treatment of the LiCoO2 and LATP powders provided the homogeneously dispersed composite granule with the granule size of several tens of micrometers. The inner dispersed LATP nanoparticles with a high conductivity of lithium ion would facilitate the ion diffusion between LiCoO2 and sulfide electrolyte during the charge and discharge reactions and increase the utilization of LiCoO2 particles inside the granule. Therefore, the solid-state battery showed capacity of 45 mAh g–1 and capacity retention of 90% after 20 cycles at 0.1 C. 3.2.9. Remarks Besides in the body structure, much more reactions and changes happen on the electrode particle/electrolyte interface and therefore appropriate surface coating and compositing can play vital role in improving the electrochemical/thermal performance mainly by stabilizing the NCM/NCA materials/electrodes. Typical surface coating with metal oxides (e.g., Al2O3 and TiO2) can protect the cathode particles from direct contact with the electrolyte and thus enhance the structural stability. Nevertheless, the metal

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oxide layers would react with the generated HF in the electrolyte and accelerate the water and HF formation during prolonged cycling. On the contrary, the utilization of metal fluoride coatings (e.g., AlF3 and LaF3) can effectively avoid the LiPF6-based electrolyte corrosion. Furtherly, a few phosphates such as AlPO4 and LaPO4 as coating shells can greatly improve the thermal stability because of the strong P=O bonds and covalency interactions between the polyanions of PO43– and metal ions. Though the enhanced interphase stability and cycling performance by the aforesaid inert coating layers of oxides, fluorides and phosphates, the surface-modified cathode materials exhibit relatively reduced reversible capacity and poor rate performance due to the poor ionic/electronic conductive shells. To facilitate the Li+/e transport and improve the electrochemical/thermal properties, either Li+ ion-conducting salt (e.g., LiAlO2 and Li3PO4) and inorganic solid-state electrolyte (e.g., LATP and LLTO) or electronic conductors (e.g., carbon) are applied as coating substances on the cathode surfaces. Furtherly, coating cathode materials with hybrid layers (e.g., dual Li+/e conductive polymers and inorganic composites) would be a promising method in the future thanks to the comprehensive advantages as both electronic and ionic conductors. Unlike the single shelled configuration, special core-shell heterostructure with double shelled layers on the cathode core is also exploited to improve the cathode performance. A lot of coating methods including solid phase-, liquid phase- and gas phase-based routes are applied for preparing the core-shelled cathode particles, however, it is still necessary to develop novel or improved coating routes to obtain complete, uniform, robust and controllable coating layers on the cathode particles at low cost and easy scale-up for their massive commercial applications. It also should be noted that the surface treatment of the cathode materials may lead to severe composition/structure changes. Albeit the improvement in the electrochemical and thermal performance, the abovementioned inert coating materials do not contribute the capacity and the thin amorphous/low-crystalline shells cannot virtually inhibit the structural change during the long-term electrochemical cycling. Therefore, other electrode materials (e.g., LiFePO4 and LiNi0.5Mn0.5O2) with better cycling or thermal stability are used to form core-shell structures for increasing the overall cathode capacity and cycling stability. To furtherly tackle the problems (e.g., core/shell separation upon lithiation/delithiation) resulted from the structural mismatch between the core and shell, combination of concentration-gradient approaches in the core/shell layers and even FCG design are developed to ensure a continuous and smooth transition in composition from the inner core to the outer shell surface for improving the structure and cycling stability. Especially, the FCG structure is suited for obtaining high-performance Ni-rich cathode particles. The synthesis procedure of the FCG particles usually contains multistep coprecipitation processes and needs to be simplified.

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In addition to the surface coating method, preparing composites with the cathode particles and multifunctional matrices offers another important structural design route to improve the electrochemical cycling performance by effectively suppressing the lithiation/delithiation-induced stress/strain and facilitating the charge transport. Particularly, the carbon-based low-dimensional nanomaterials and solid-state electrolytes show promising application prospects. 3.3. Morphology design Aside from the modifications on the composition and structure levels, controlling the morphology characteristics of the cathode particles offer another important materials design route for improving the cycling performance, thermal stability, tap/packing density, etc. 3.3.1. Special morphologies (size/pore/shape/configuration control) The electrode reactions occur at the surface and require the transportation of electrons/ions, so the structures/morphologies of the electrode materials play an important role in the electrochemical properties. Small particles are desired because of the effective alleviation of the volume change-induced stress/strain, and the large specific surface area and short diffusion distance [3, 502-505]; however, too small particles induce more surface reactions and the high reactivity of the nanoparticles could be disadvantageous in terms of safety and stability during the long operational lifetimes [422, 506-510]. Therefore, proper particle sizes sometimes possess the best performance. Zhang et al. [511] synthesized one-dimensional (1D) LiNi0.8Co0.15Al0.05O2 micro-rods by two-step co-precipitation method using mixed metal oxalate micro-rods as sacrificial templates. Because of the special 1D structure, the 1D NCA exhibited higher capacities of 168 and 147 mAh g–1 and retention rates of 93% and 72% than the common NCA (133 and 75 mAh g–1 and 87% and 41%) after 100 cycles at 1 C within 2.7–4.3 V at 25 and 55 oC, respectively. Luo et al. [512] fabricated hexagonal nanosheets of LiMn0.075Co0.775Ni0.15O2 with ~85 nm thickness and exposing (101) facets via coprecipitation and subsequent heat treatment. The nanostructure methodology can resolve the kinetic problems of Li+ ions and electrons in electrodes, and thus the electrode exhibited high capacities of 151.3 and 135.9 mAh g–1 and capacity retentions of 85.1% and 84.7% after 50 cycles at 300 and 3000 mAh g–1 under a high cut-off voltage of 4.4 V, respectively. Cao et al. [513] synthesized various holey 2D ultrathin cathode nanosheets (LiNi0.8Co0.15Al0.05O2, LiNi0.8Co0.1Mn0.1O2, LiNi0.6Co0.2Mn0.2O2, and LiNi0.5Co0.2Mn0.3O2) through a combined microwave-assisted and solid-state reaction and investigated their cycling performance. The NCA, NCM811, NCM622 and NCM523 electrodes showed high initial capacities of 205.4, 198.6, 184.5, and 168.6 mAh g-1, and maintained ~90% capacity retention and ~100% Coulombic efficiency after 50 cycles at 0.1 C between 2.7–4.3 V. The electrodes also showed capacities of 169.3, 162, 149.6 and 134.1 mAh g-1 at 5 C, respectively, and maintained ~90% of the

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5C capacities at 10 C. The improved cycling performance were attributed to the unique 2D nanosheet constructure with abundant holes, which can reduce the Li+ ion diffusion distance and offer more electrochemical active sites. A few cathode materials with porous, hollow and tubular shapes also perform well during electrochemical cycling, due to the increased surface area and effective alleviation of the volume expansion/contraction during lithiation/delithiation and inhibition of the micro-cracking and pulverization [424, 514-522]. Chen et al. [523] prepared porous LiNi1/3Co1/3Mn1/3O2 microspheres through a two-step route using fluffy MnO2 as sacrificial template. This structure can effectively increase the contact area with electrolyte, shorten Li+ ion diffusion path and improve the Li+ ion mobility. As a consequence, the porous electrode possessed higher capacity of 179 mAh g–1 and retention rate of 90% than the bulk electrode (88 mAh g–1 and 50%) after 50 cycles at 0.1 C in a potential range of 2.5–4.5 V. Moreover, the porous electrode showed much higher rate capacities of 208, 178, 164, 149 and 120 mAh g–1 than the bulk electrode (182, 123, 88, 61 and 24 mAh g–1) at 0.1, 0.5, 1, 2 and 5 C, respectively. Zou et al. [524] prepared multi-shelled Ni-rich LiNixCoyMnzO2 (x=0.8, 0.7, 0.65 and 0.5) hollow fibers through cationic ion exchange and following calcination using alginate fibers as sacrificial templates (Fig. 23A). The hollow fibers of ~10 μm in diameter are highly porous and had double or triple shells of hundreds of nanometers in thickness with wide void space between the shells, which would effectively increase the electrode/electrolyte contact area and enhance the diffusion rates for both Li+ ions and electrons (Fig. 23B-C). When increasing the Ni content from 0.5 to 0.8, the Li/Ni cation mixing/exchange ratio increased from 0.0485 to 0.0675, which was lower than the value of 0.085 for bulk LiNi0.65Co0.25Mn0.1O2. The initial discharge capacities were 229.9, 217.2, 211.5 and 204.4 mAh g–1 at 0.1 C between 2.5–4.5 V, and the electrodes still remained high capacities of 200.4, 197.0, 193.1 and 190.5 mAh g–1 and capacity retentions of 91.2%, 91.9%, 92.2% and 94.2% after 200 cycles at 0.5 C, for x=0.8, 0.7, 0.65 and 0.5, respectively (Fig. 23D). Owing to the special morphology characteristics, the multi-shelled LiNi0.8Co0.1Mn0.1O2 electrode also delivered stable rate capacities of 207.4, 196.6, 183.4 and 172.7 mAh g–1 at 1, 2, 5 and 10 C, respectively (Fig. 23E).

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Fig. 23 (A) Schematic illustration of the preparation of multi-shelled NCM hollow fibers by (a) cationic ion exchange in the alginate fibers and (b) following calcination, (B) SEM and (C) TEM images of the LiNixCoyMnzO2 (x=0.5) fibers, and (D) cycling performance at 0.5 C between 2.5–4.5 V and (E) rate capacities of the NCM fibers. Adapted with permission [524]. Copyright 2016 WileyVCH.

Controlling the size and porosity allows for easy access of electrolyte to the materials surface, and nanoporous microparticles have been reported to exhibit better cycling performance [15, 522, 525-528]. Wu et al. [529] compared the cycling performance of Li1.2Mn0.54Ni0.13Co0.13O2 nanoparticles and secondary microspheres constituted with primary nanoparticles synthesized by sol-gel and co-precipitation, respectively. The nanoparticle electrode showed higher initial capacities of 242 and 216 mAh g–1 than the microparticle electrode (226 and 187 mAh g–1) at 0.5 and 2 C, respectively, due to the shorter Li-ion diffusion length in the nanoparticles. However, the microparticle electrode exhibited higher capacities of 203 and 165 mAh g–1 and retention rates of 90% and 88% than the nanoparticle electrode (181 and 122 mAh g–1, 75% and 56%) after 60 cycles at 0.5 and 2 C, respectively, because the hierarchical configuration can effectively hinder the crack growth during cycling. Chen et al. [530] synthesized uniform and dense LiNi0.9Co0.07Al0.03O2 microspheres with well-assembled nanoparticles and low degree of Li+/Ni2+ mixing by optimizing the co-precipitation and calcination conditions. Due to the special heterostructure and low mixing degree, the NCA sample with a tap density of 2.52 g cm–3 showed high cycling stability at various temperature (capacity retentions of 93.2% and 83.8% after 100 cycles at 1 C at 25 and 55 oC, respectively) and excellent thermal stability (heat generation of 517.5 J g–1 delithiated at 4.3 V). Moreover, the coin-type NCA/graphite full-cell with the cathode mass of ~6 mg showed high capacity of 166.6 mAh g–1 and retention rate of 91.7% after 100 cycles at 1 C within 2.0–4.3 V. Correspondingly, this configuration provided the energy densities of 619.8 and 278.9 Wh kg–1 for the NCA cathode material and the

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18650 cell with ~45 wt% cathode mass loading based on the average working potential of 3.72 V, respectively. The shape of cathode particles holds the key in determining the battery performance by affecting the mobility of charge carriers, electrochemical reactions and the lattice strain during lithiation/delithiation [531, 532]. Jagannathan et al. [533] fabricated LiNi1/3Mn1/3Co1/3O2 particles using a mixed hydroxyl-carbonate precursor, and the sample with a faceted, edge-blunted polyhedral shape showed a higher initial capacity of 187 mAh g–1 than the spherical particles at 0.1 C and a higher capacity of 153 mAh g–1 with a capacity retention of 82% after 50 cycles, mainly attributed to the unique morphology which facilitated the charge transfer. In particular, a unique particle morphology was developed with the nanorod primary particles radially aligned toward the outer shells of the secondary microsphere particles for significantly enhancing the cycling retention, since this can reduce the number and size of the voids in which the electrolyte can fill, minimize the interface area between the active materials and the electrolyte, and hinder the side reactions [40, 471, 483]. Not only the size but also the shape (e.g., spheres and ellipse) of the cathode particles can determine the tap/compact density and affect the battery performance [36, 527, 534, 535]. Meanwhile, the electrode materials with consistent shape and optimal size could ensure uniform reversible insertion/extraction during repeated chargedischarge processes [535]. Liu et al. [534] studied the cycling performance of irregular Li1.2Mn0.54Ni0.13Co0.13O2 secondary particles composed of spherical or quasi-spherical primary nanoparticles with small pores between them and spherical secondary particles composed of pillar-shaped primary nanoparticles with larger gaps between them. Because of the larger surface area and more active crystal planes of the irregular particles, the cathodes with irregular particles showed higher tap density of 2.17 g cm–3, gravimetric capacity of 259 mAh g–1 and volumetric capacity of 926 mAh cm–3 than the spherical particle-based electrode (2.03 g cm–3, 232 mAh g–1, and 791 mAh cm–3). The irregular particle-based electrode also exhibited a higher capacity of ~137 mAh g–1 than the spherical particle-based electrode (~125 mAh g–1) after 100 cycles at 250 mA g–1. 3.3.2. Single-crystal cathode particle We all know that most of the layered NCM/NCA materials are in the form of micro-sized secondary particles with the agglomeration of nanosized (or sub-microsized) primary particles, and plenty of voids formed between the primary particles result in the low tap and compression/packing densities (≤3.6 g cm–3 in electrode density) and volumetric energy density (≤700 mAh cm–3) [48]. For example, the commercial LiNi0.6Co0.2Mn0.2O2 particles exhibit relatively low electrode density of ~3.4 g cm–3 and volumetric energy density of ~2260 Wh cm–3, compared to the commercial singlecrystalline LiCoO2 particles (~3.8 g cm–3 and ~2300 Wh cm–3). Besides, the heterogeneous configuration could be easily destroyed during the electrode pressing

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process ascribed to the collapse of the aggregated secondary particles, leading to more side reactions with the electrolyte and phase transition due to the more exposure of the particle surface areas and highly instable Ni3+ ions to the electrolyte [48]. Although elevating the compression force in the electrode fabrication process can increase the electrode density because of the decrease in the porosity, it will reduce the electrolyte wettability to the cathode particles attributed to the blockage of the electrolyte (LiPF6) pathway into the electrode, resulting in the poor rate capability and thermal stability [536-539]. Moreover, the thermal decomposition is greatly facilitated in the highdensity cathodes, because the high electrode density hinders the decomposition reactions of LiPF6 and subsequent formation of polymeric compounds that could suppress the oxygen release from the cathode particle surfaces [539-541]. Additionally, the volume change induced by the electrochemical reactions during the cycling process also accelerates the crackle and pulverization of the cathode particles, and thus leads to the severe capacity fading [17, 34, 542-545]. Inspired by the excellent cycling performance of the single-crystal LiCoO2 particles, researchers begin to develop single-crystal NCM/NCA particles for enhancing the stability of the crystal structure, increasing the tap/packing and energy densities, and improving the mechanical processing property. The single-crystalline NCM/NCA particles with lower surface area also have higher safety in terms of thermal stability than the secondary samples, since the morphology robustness can suppress the crack evolution, prevent the side reactions with the electrolyte and hinder the gas release [132, 546]. High-temperature treatment of the metal precursors (usually in the form of singlecrystal structure other than polycrystalline structure) synthesized by modified coprecipitation, sol-gel and spray-drying offers an simple route for obtaining single-crystal particles; however, it usually results in coalescence of the particles into irregular secondary particles and non-uniformity in the shape and size [36, 90, 92, 98, 546-549]. In addition, the high-temperature sintering process may also lead to over-growth of the partial particles with diameters greater than 10 μm [36]. Lin et al. [90] synthesized single-crystal LiNi1/3Co1/3Mn1/3O2 particles of several hundreds of nanometers in diameter by an improved co-precipitation approach by employing H2O2 as an additive followed by ozone-assisted thermal treatment. Notably, H2O2 acted as oxidant and dispersant during the co-precipitation process for oxidizing the metal ions and suppressing the agglomeration of the precursors by giving out O2, whereas O3 could further oxidize NCM to enhance the structure stability during the calcination process. Consequently, the single-crystal NCM electrode showed a high initial capacity of 209 mAh g–1 at 0.1 C (28 mA g–1) and maintained a capacity of 176 mAh g–1 with a retention rate of 84% after 50 cycles in a voltage range of 2.5–4.6 V. The single-crystal electrode also showed high capacity of 156 mAh g–1 and capacity retention of 82% after

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50 cycles at 1 C. Wang et al. [548] prepared single-crystal LiNi0.6Co0.2Mn0.2O2 particles by a precipitation-hydrothermal reaction and subsequent calcination. When the calcination temperature increased from 750 to 850 oC, the average size of the singlecrystal particles increased from 300 to 800 nm; however, when the temperature increased to 900 oC, the particles growed very fast with a wide size distribution and aggregated more severely. The increase of the annealing temperature can also result in the crystallization improvement of the cathode particles and reduce the cation mixing. The single-crystal structure can be beneficial to the transportation of Li+ ions along the crystal, and meanwhile the sub-micron sizes would reduce the side reactions between the particles and the electrolyte due to the relatively low surface areas. As a result, the single-crystal sample displayed high initial capacities of 172.7 and 153.6 mAh g–1 and capacity retentions of 93.5% and 80.5% after 100 cycles in a voltage range of 2.8–4.3 V at 0.5 and 10 C, respectively. When cycled within 2.8–4.6 V, the sample showed initial capacities of 191.6 and 167.0 mAh g–1 and capacity retentions of 78.6% and 66.0% after 50 cycles at 0.5 and 10 C, respectively. A molten salt-assisted flux growth process is also exploited for synthesizing singlecrystal cathode particles by heating both the secondary particles of the co-precipitated cathode precursors and other chemical reagents (e.g., LiCl, LiNO3, Li2SO4, NaCl, Na2SO4, KCl, Li2MoO4 and Na2MoO4) at relatively low temperatures [36, 546, 550554]. The molten salts of the chemical reagents can lower the melting temperatures of the secondary precursors, and the cathode materials would dissolve and re-precipitate at the surfaces for growing single-crystal cathode particles with narrow size distribution [555]. Moreover, adjusting the type and concentration of the chemical reagents can also affect the size and shape of the single-crystal particles. Nevertheless, owing to the unavoidable side reactions with the chemical reagents during the annealing process, this synthesis approach could result in the formation of solid solution phases in the bulk structure and impurities on the cathode particle surfaces, which need to be removed by an additional washing process [36]. Kimijima et al. [36] prepared individually-dispersed single-crystal LiNi1/3Co1/3Mn1/3O2 particles using a molten Li2MoO4-assisted sintering process. When increasing the solute concentration from 20 to 80 mol% with a sintering temperature of 1000 oC, the NCM particle size increased from 2.3 to 4.4 μm (Fig. 24AC); when increasing the calcination temperature from 700 to 900 oC, the particle size increased from 0.09 to 1.1 μm with more individual dispersity. Because of the side reactions between the molten Li2MoO4 and the cathode materials during the calcination process, amorphous layers formed on the NCM particle surfaces (Fig. 24D). Thus, an additional washing process with warm water was utilized to remove the impurities on the particle surfaces, which resulted in the formation of lithium deficiency in the NCM particles. In order to recover the lithium deficiency, the NCM particles were heated again at 800 oC for 10 hours with the addition of LiOH (Fig. 24E). Because of the

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removal of the amorphous layers and recovery of lithium occupancy, the reheated sample showed a higher initial capacity of 153 mAh g–1 than the pristine sample (132 mAh g–1) at 0.1 C in a voltage range of 2.8–4.3 V (Fig. 24F). The crystals grown at 800 oC with post-heat-treatment (NCM-800HT) represented the best rate performance, and the capacity at 2000 mA g–1 was much higher than that of the commercial sample (Fig. 24G). Moreover, the reheated samples at a higher calcination temperature of 1000 oC displayed higher capacity of 102 mAh g–1 and capacity retention of 78% than the reheated samples at lower calcination temperatures of 900 (100 mAh g–1 and 75%) and 800 oC (95 mAh g–1 and 68%) after 100 cycles at 1 C (Fig. 24H), implying that the smaller particles with larger surface areas could cause higher capacity fading due to the more severe side reactions.

Fig. 24 SEM images of the pristine NCM particles calcinated at 1000 oC with the co-precipitated solution concentrations of (A) 80, (B) 40 and (C) 20 mol%, respectively, TEM images and corresponding SAED patterns of the (D) pristine and (E) reheated NCM samples with a calcination temperature of 900 oC (denoted as NCM-900 and NCM-900HT, respectively), (F) capacity-voltage profiles of (a) NCM-900 and (b) NCM-900HT at 20 mA g–1 between 2.8–4.3 V, and (G) rate capacities and (H) cycling performance at 200 mA g–1 of the commercial (NCM-T), conventionallygrown (NCM-N), pristine (NCM-900) and reheated (NCM-800HT, NCM-900HT and NCM1000HT) NCM samples. Adapted with permission [36]. Copyright 2016 The Royal Society of Chemistry.

3.3.3. Remarks The morphological characteristics (size/pore/shape/configuration) of the cathode materials not only affect the structure stability, charge transfer and electrochemical/thermal performance but also the tap/electrode/energy densities. Smallsized (especially nano-sized) and porous/hollow/tubular cathode particles exhibit high morphology stability and rate performance due to the high surface-to-volume and effective inhibition of the volume variation during lithiation/delithiation processes, but the much more surface side reactions and low tap density are detrimental to their

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practical applications. Thus, it is necessary to control the size, porosity and shape of the cathode particles. Recently, single-crystal cathode particles are becoming a research hotspot due to the advantages of the structural/morphology stability, electrode fabrication process and tap/energy densities. However, the current synthesis methods such as high-temperature treatment and molten salt-assisted flux growth are far from the commercial viability, because of the agglomeration of the particles, non-uniformity in shape and size, formation of solid-solution phases in the body structure, impurities on the particle surfaces, complex processing, difficulty in scale-up, etc. 3.4. Co-modification The individual composition, structure and morphology modification on the NCM/NCA materials can affect the properties of the cathode materials, and therefore it is reasonable to completely and greatly improve the battery performance by appropriate integration of these modification methods on the basis of the synergistic effect. Amongst the co-modification approaches, the combination of body doping and surface coating is usually used to enhance the electrochemical and thermal properties of the cathode materials by meanwhile stabilizing both the lattice structure and the electrode/electrolyte interface [556, 557]. A few metal elements (e.g., Mg, Nb, Zr and Ti) may diffuse into the bulk structure during the heating process, and therefore it is easy to obtain co-modified cathode materials with both the lattice doped and surface coated by metals and the metal-based compounds, respectively [210, 288, 325, 558-562]. Huang et al. [374] prepared Zrmodified LiNi1/3Co1/3Mn1/3O2 particles (size: ~8 μm, tap density: 2.4 g cm–3) through modified co-precipitation and subsequent heat treating. During the high-temperature process, a part of Zr covered the surface of LiNi1/3Co1/3Mn1/3O2 to form a Li2ZrO3 coating layer, while the other part diffused into the LiNi1/3Co1/3Mn1/3O2 bulk to modify the lattice structure (the lattice parameters varied with the Zr content). 1 wt% Zrmodified LiNi1/3Co1/3Mn1/3O2 exhibited higher capacity of 165 mAh g–1 and capacity retention of 99% after 100 cycles at 0.5 C in a potential range of 3.0–4.5 V than the Zrfree sample (132 mAh g–1 and 78%). In addition, the capacity of the bare material decreased faster than that of the 1 wt% Zr-modified material when the discharging current increased from 0.2 to 5 C. The improved high-voltage cycling performance can be ascribed to: (1) the Li2ZrO3 on the surface of the LiNi1/3Co1/3Mn1/3O2 particles acted as a protective layer to reduce the direct contact between the electrode and electrolyte, and thus suppressed the dissolution of the transition metal elements and electrolyte decomposition; (2) the strong Zr–O bond of Li2ZrO3 on the electrode surface was helpful for lowering the activity of oxygen on the electrode surface at high voltage; (3) the Zr that diffused into the crystal lattice stabilized the structure during cycling. Wang et al. [325] prepared LiNi0.8Co0.15Al0.05O2 particles with both Ti substitution in Ni sites in the TmO2 slabs and 22 nm heterogeneous TiO2 coating layers on the particle surfaces

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for reducing the cation mixing degree and protecting the particle from acid attacking, and thus the modified NCA electrode showed high capacity of 182.4 mAh g–1 and retention rate of 85.0% after 200 cycles at 1 C. Likewise, cathode materials can be body doped with single or two elements and meanwhile surface coated with other substances. Wang et al. [563] prepared Zr-doped LiNi0.5Co0.2Mn0.3O2 particles with the surfaces coated by conductive polypyrrole (PPy) layers (NCM-Zr/PPy) via solid-state reaction followed by chemical oxidation polymerization. When cycled at 2 C for 200 cycles in a potential range of 3.0–4.6 V, the NCM-Zr/PPy electrode displayed a higher capacity of 138 mAh g−1 and more stable stability with a higher capacity retention of 78.8% than the NCM (100 mAh g−1, 56.8%) and NCM-Zr (118 mAh g−1, 67.0%) electrodes. The NCM-Zr/PPy electrode also showed a discharge capacity of 136 mAh g−1 at 10 C, which was much higher than those of the NCM and NCM-Zr cathodes (96 and 116 mAh g−1). Furthermore, the NCM-Zr/PPy particles virtually maintained the spherical shape even after the long-term cycling, while the NCM particles suffered from severe damage to collapse into small fragments and only a part of the NCM-Zr particles still preserved the original spherical morphology. The improvement in the cycling stability are attributed to the synergistic effect of the bulk Zr doping and surface PPy coating: (1) the PPy thin film on the particle surface established a conductive network and accelerated the electronic conductivity; (2) the coated PPy layer avoided the direct exposure of the NCM particles to the electrolyte and hindered the continuous HF erosion from the decomposition of electrolyte; (3) the Zr substitution enlarged the Li slab distance and was beneficial to Li+ ion transportation. The combination of doping and FCG offers another important co-modification approach [564, 565]. The Sun group very recently prepared 1–2 mol% Al-doped LiNi0.76Co0.09Mn0.15O2 (average composition) particles with FCGs of Ni, Co and Mn spanning the particle core to the surface (Fig. 25A-D) [566]. Both the lattice unit cell volume and the cation mixing ratio decreased with the increase of the Al content in the cathode particles. As the Al content increased to 2 mol%, the initial discharge capacity slightly decreased from 208 to 203 mAh g−1 at 0.1 C, whereas the capacity retention increased from 93% to 95% after 100 cycles at 0.5 C in a voltage range of 2.7–4.3 V. The pouch-type FCG-LiNi0.76Co0.09Mn0.15O2/graphite full-cell showed a higher capacity retention of 87.9% than the commercial LiNi0.82Co0.14Al0.04O2/graphite full-cell (79.6%) after 1000 cycles at 1 C in a voltage range 3.0–4.2 V, because of the special FCG constructure; however, the 1 mol% and 2 mol% Al-doped LiNi0.76Co0.09Mn0.15O2/graphite full-cells exhibited higher capacity retentions of 93.7% and 95.0%, respectively (Fig. 25E), because the combination of the Al-doping and FCG design can synergistically enhance the structural and surface stabilization as well as thermal stability (Fig. 25F-G). After the long-term cycling, the LiNi0.82Co0.14Al0.04O2

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particles were completely pulverized, whereas the Al-doped FCG particles kept intact without visible cracks, further evincing the improved mechanical integrity by the Al doping and FCG construction (Fig. 25H-K).

Fig. 25 (A) SEM images of FCG-LiNi0.76Co0.09Mn0.15O2 (denoted as FCG76) particles, (B) brightfield TEM image of 1 mol% Al-doped LiNi0.76Co0.09Mn0.15O2 (Al-1 FCG76) primary particles, (C) EPMA composition profile of a Al-1 FCG76 secondary particle, (D) XRD spectra of the different FCG76 powders, (E) extended cycling performance of LiNi0.82Co0.14Al0.04O2 (NCA82) and FCG76 samples in pouch-type full-cells with graphite anodes at 1 C, HR-TEM images and corresponding SAED spectra and FFT patterns of the (F) FCG76 and (G) Al-1 FCG76 samples after 1000 cycles, and SEM images of the cycled (H) NCA82, (I) FCG76, (J) Al-1 FCG76 and (K) Al-2 FCG76 samples. Adapted with permission [566]. Copyright 2017 American Chemical Society.

4. Electrode design Apart from the materials design on the composition, structure and morphology of the layered NCM/NCA cathode materials, proper electrode design on the electrode composition, fabrication process and environment condition [567, 568] (Fig. 1 and 26) is also important to facilitating the large-scale commercial applications of the layered NCM/NCA materials, mainly attributed to the improvements in the structure stability and charge transport.

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Fig. 26 Electrode design methods.

4.1. Composition adjustion In addition to the active electrode materials, the additives such as polymer binders and conductive graphite materials with small amounts also greatly affect the electrode structure, porosity, Li+/e‒ conductivities, mechanic strength, thermal stability, battery performance, etc [569-571]. For instance, Manthiram group [572] recently found that the carbon additive-driven interphase formed in situ on LiNi0.7Co0.15Mn0.15O2 particles before the cell operation served as a passivation film against the deleterious interfacial reactions with the electrolyte with the assistance of time-of-flight secondary-ion mass spectrometry (TOF-SIMS) technology. Ideal positive electrodes can also be obtained by optimizing the composition (e.g., polymer binder and conducting additive) and content. 4.1.1. Conductive additives Graphite materials such as carbon black (CB) and acetylene black (AB) are usually used as conductive additives in commercial Li-ion batteries, because of their advantages of high electric conductivity, high thermal conductivity, low weight-to-volume ratio, high resistance to acid and alkali, low cost, environmental friendliness, etc. The optimization of graphite additive and the electrode porosity are needed to obtain both high electric conductivity and fast Li+ ion diffusivity. Moreover, the graphite crystallization degree, shape, size, porosity, purity, and surficial defects (e.g., oxygencontaining groups) of the graphite additives play an import role in determining the battery cycling performance. Qi et al. [573] investigated the effect of graphite crystallization of CB (average size of 31 nm) on the battery cycling performance, and found that the high-crystallization CB led to better cycling stability of the LiNi0.5Mn1.5O4/Li cells, mainly owing to the removal of the functional oxygen-

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containing groups and the inhibition of the side reactions with the electrolytes. In another work, Lee et al. [492] used nitric acid to treat Super P carbon black particles to change the particle surface from hydrophobic to hydrophilic for better dispersing the cathode primary particles in water, and the resulted electrode with uniformly-distributed carbon black particles can facilitate the charge transfer and thus showed greatlyimproved rate capability and cycling stability. Maeyoshi et al. [574] furtherly studied the effect of specific surface area of acetylene black (AB) and ketjen black additives on the LiCoPO4/Li cell performance. The high-surface-area (853 m2/g) ketjen black-based cells showed the highest initial capacity, but the capacity degraded fast due to the much more side reactions; while the low-surface-area (39 m2/g) AB-based cells showed the low capacity due to the insufficient contact between the cathode particles and AB. In sharp contrast, the medium-surface-area (68 m2/g) AB led to the highest capacity of 70.5 mAh g−1 and capacity retention of 56.6% after 100 cycles at 0.1 C within 3.0‒5.1 V, ascribed to the effective conductive networks in the cathodes and the suppression of irreversible decomposition reactions of the electrolyte and the anion interactions. Carbon fibers are also used as conductive additives, due to the low density (usually ~0.05 g/cm3), high bulk conductivity and high thermal conductivity. The onedimensional fibers can also strengthen the flexibility and mechanical stability of the resulted electrodes [575]. Nevertheless, the complicated synthesis process and high cost hinder their applications in Li-ion batteries. Besides, the fibers with tubular shape cannot entirely cover the active particles like the nano-sized carbon blacks, and the electron transport across the active particle surface is also restrained. To overcome this problem, the combination of the carbon fiber and carbon black seems a good choice by integrating their synergic advantages [576]. Bian et al. [577] compared the battery performance based on the carbon nanofiber/CB mixture additives with various contents and found that the Li1.18Ni0.15Co0.15Mn0.52O2/Li cells with the 3 wt% carbon fibers and 12 wt% CB showed the highest capacity of 239 mAh g−1 and capacity retention of 94.4% after 50 cycles at 40 mA g−1 within 2.0‒4.8 V. Other carbon nanomaterials such as carbon nanotubes (CNTs) and graphene are also usually utilized to replace the traditional conductive additives, due to their high electrical conductivity, high flexibility and the cross-linked charge-transfer paths in the electrodes [578-581]. The commonly-used binding and conducting agents do not contribute the capacity, and their high content (usually 2–10 wt%) in the cathodes results in decreased capacity; however, the less use of the carbon-based nanomaterials (0.5–2 wt%) with enhanced conductivity in the electrode can improve the energy and power densities [582]. Alberto et al. [582] studied the effect of multi-walled carbon nanotubes (MWCNTs) for partial substitution of carbon black as conductive addictive in LiNi1/3Co1/3Mn1/3O2 particle-based electrode. The addition of MWCNTs significantly enhanced the rate capability of the LiNi1/3Co1/3Mn1/3O2 cathodes at all investigated

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charge-discharge rates of 0.25–5 C (e.g., 87 and 58 mAh g–1 at 5 C for the MWCNTand carbon black-based electrodes, respectively). Varzi et al. [583] investigated the influence of pristine and carboxylated MWCNTs on the cell performance of NCM and LiFePO4 electrodes. The carboxylated MWCNTs can effectively prohibit the agglomeration of the electrode materials and improve their connection, and thus the LiFePO4 electrodes with the carboxylated MWCNTs showed the best cycling performance. However, the NCM electrodes with the carboxylated MWCNTs exhibited the poorest performance, attributed to the inhibition of the Li+ ion insertion/extraction by the low-conductivity carboxylated MWCNT layers on the active particles. Du et al. [584] furtherly prepared LiNi0.5Co0.2Mn0.3O2/graphite pouch cells using the CNT and CB additives and compared the battery performance and cost. The CNTs can evenly disperse in the cathodes, uniformly cover the cathode particles and form threedimensionally conductive networks (Fig. 27A-C), and improve the cathode conductivity compared to the CB-based cathodes (Fig. 27D). As a result, the CNT-based cells showed higher capacity retention of 99.4% and lower voltage polarization of 130 mV than the CB-based cells (94.6% and 440 mV) after 200 cycles at C/3 within 2.5‒4.2 V (Fig. 27E). Moreover, the NCM loading increased from 90 wt% (5 wt% CB and 5 wt% binder) to 96 wt% (0.3 wt% CB, 1.2 wt% CNT and 2.5 wt% binder) with the energy density increase by 11.7%, and the cost saving of 33% and 58% at 2 C and 3 C, respectively (Fig. 27F).

Fig. 27 SEM images of (A) CB electrode (90 wt% NCM, 5 wt% CB, and 5 wt% binder), (B) CNT-A electrode (92.5 wt% NCM, 0.7 wt% CB, 1.8 wt% CNT, and 5 wt% binder), and (C) CNT-B electrode (96 wt% NCM, 0.3 wt% CB, 1.2 wt% CNT, and 2.5 wt% binder) after 200 cycles, (D) cell resistance with different positive electrodes at C/3, (E) cycling performance of the different electrodes, and (F) calculated cathode energy cost of three electrodes at C/3, 1 C, 2 C and 3 C rate. Adapted with permission [584]. Copyright 2018 Elsevier.

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The effects of mesoporous graphene and various graphite additives of carbon black (CB) and conductive graphite (CG) on the LiNi1/3Co1/3Mn1/3O2/Li cell performance were investigated by Liu et al [585]. It was found that the graphene/CB-containing cells showed the best rate capacity of 78 mAh g−1 at 2000 mA g−1 within 2.5‒4.1 V, due to the reduced electrode resistance by the high-conductivity graphene additive. Liu et al. [586] investigated the effect of graphene size on the LiFePO4/Li cell performance, and found that the small-size graphene additive (1.0‒2.0 µm) resulted in highest capacity and rate performance (1 wt% graphene and 9 wt% Super P as conductive binder). Shi et al. [587] compared the impacts of three graphene additives synthesized by intercalation exfoliation, electrolysis exfoliation and modified Hummers methods on the battery cycling performance, and found that the graphene with larger size, high crystallinity and high conductivity cannot result in the best battery performance. However, the graphene with proper sheet size and moderate oxygen-containing groups and electrical conductivity can greatly improve the cycling performance. That is, the size, surface chemistry and electrical conductivity of the graphene additives are essential factors affecting the battery performance. The work by Juarez-Yescas et al. [588] also proved the importance of the oxygen-containing groups and electrical conductivity of graphene to the battery cycling performance. 4.1.2. Binders Binders are indispensable components in Li-ion batteries, and play an important role in holding the electrode materials together and ensuring a superior bonding of the electrodes to the current collectors [589]. Because of the combined advantages such as wide electrochemical window, high adhesivity with electrode materials, high viscosity of the organic slurry, low swelling ratio and dissolubility in electrolyte, and high flexibility, polyvinylidene difluoride (PVDF) is usually used as binder for both cathodes and anodes. However, PVDF binder is much expensive, and the use of organic solvents such as the toxic and low-volatility N-methyl-2-pyrrolidone (NMP) is detrimental to the environment and battery manufacturing. Moreover, the PVDF binder is nonconductive to Li+/e‒. Therefore, novel binders are needed to replace the PVDF binder for improving the cycling performance and safety of the NCM/NCA-based cells and also decreasing the manufacturing cost. Conjugated polymers including poly(3,4-ethylenedioxythiophene) (PEDOT), polypyrrole (PPy), poly(3-hexylthiophene-2,5-diyl) (P3HT) and polythiophene (PT) are recently used to replace the insulating PVDF binder, because of their high electroconductivity, electrochemical stability, and easy integration into electrode structures. Dunn et al. [590] prepared LiNi0.8Co0.15Al0.05O2-based electrodes with P3HT and CNT as conductive binder and additive, respectively. The oxidization of the P3HT during the charge-discharge potential range endowed the high conductivity. In addition, the P3HTCNT binder system can inhibit the PF6– decomposition and the generation of HF,

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suppress the formation of SEI film, and prohibit the inter-granular cracking in the NCA secondary particles. As a result, the PSHT-CNT modified electrode showed improved capacity of 80 mAh g–1 (80% of the initial capacity) after 1000 cycles at 16 C, a value that was four times greater than that obtained in the PVDF-based electrode. A few polymers such as polyamide-imide (PAI) and polyimide (PI) with high glasstransition temperatures can effectively enhance the thermal stability of the cathodes and safety of the cells. A study by Morishita et al. [591] showed that 5 wt% PAI-containing LiNi1/3Co1/3Mn1/3O2/SiO full cells displayed higher capacity of 88 mAh g−1 and capacity retention of 80% than the 5 wt% PVDF-containing cells (25 mAh g−1 and 23%) after 500 cycles at 5 C and 60 oC within 1.5‒4.15 V, because of the higher adhesion and thermal stability of the PAI binder than the PVDF binder, and the effective suppression of the electrolyte decomposition by the adhesive PAI on the NCM surface. Fluorinated polyimide (FPI) with abundant –CF3 and –COOH groups can improve the thermal and electrochemical stability, and meanwhile form on the cathode particle surface as a protection layer, and thus the Li1.2Ni0.132Co0.13Mn0.54O2/Li cell displayed higher capacity of 223‒198 mAh g−1 and capacity retention of 89% than the PVDF-based cells after 100 cycles at 0.2 C and 55 oC within 2.5‒4.7 V [592]. Qian et al. [593] found that PI binder better adhered and wrapped the cathode particles than PVDF binder (Fig. 28A-B), and also showed less solubility and swelling in the commercial electrolyte (Fig. 28C-D). The PI binder-based LiNi1/3Co1/3Mn1/3O2/graphite pouch cells also showed comparable capacity retention of 91.4% after 500 cycles at 0.5 C and rate capability within 2.75‒4.2 V (Fig. 28E). Furthermore, the PI-based pouch cells passed the overcharge (up to 10 V) safety test, but the PVDF-based cells easily took fire during the safety test (Fig. 28F-G).

Fig. 28 SEM images of the (A) PVDF- and (B) PI-containing LiNi1/3Co1/3Mn1/3O2 cathodes, the swelling comparison of PVDF and PI in 1 M LiPF6:EC/DMC/EMC electrolyte for (C) 0 and (D) 48 h, (E) rate performance of the PVDF- and PI-containing LiNi1/3Co1/3Mn1/3O2/graphite pouch cells

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within 2.75‒4.2 V, and voltage and temperature profiles of the (F) PVDF- and (G) PI-containing pouch cells at a 1 C rate overcharge. Adapted with permission [593]. Copyright 2016 Elsevier.

Water-soluble binders such as carboxymethyl cellulose (CMC) [594, 595], alginate [596, 597], guar gum (GG) [598], and polyacrylic latex (PAL) [599] are recently used for both anodes and cathodes [600], owing to their advantages such as low cost, environmental friendliness and simplification of the battery manufacturing. A comparative study by Xu et al. [601] disclosed that the CMC-based LiNi1/3Co1/3Mn1/3O2/Li cathodes had higher capacity of 141.9 mAh g−1 and capacity retention of 90.1% than the alginate- and PVDF-based electrodes (126 mAh g−1 and 89.2%, and 111.7 mAh g−1 and 86.3%) after 200 cycles at 0.5 C within 2.5‒4.6 V and also better rate performance, attributed to the lower charge transfer resistance of the high-ion-conductivity CMC binder-based electrodes. Though the obvious advantages of the water-soluble binders, the strong reactivity between the metal oxide materials and water results in the leaching of Li+ and metal ions from the cathodes and the increase of pH that in turn causes the corrosion of the aluminum current collectors. To suppress the corrosion of Al current collectors, Loeffler et al. [602] used a mixture binder of polyurethane (PU) and CMC. Because of the protection layers of PU on the Al current collectors and LiNi1/3Co1/3Mn1/3O2 particles, the 1 wt% CMC/2 wt% PU-containing NCM/Li half cells showed higher capacity of 120 mAh g−1 than the 2 wt% CMC/1 wt% PU- and 3 wt% CMC-containing cells (112 and 104 mAh g−1) after 100 cycles at 1 C within 3.0‒4.3 V. 4.1.3. Blending with other cathode materials In contrast to other strategies such as doping, coating and compositing, blending with other cathode materials is a more convenient method for eliminating the irreversible capacity loss of the layered NCM/NCA cathode materials [603-606]. Particularly the Ni-rich and Li-rich layered cathode materials have large irreversible capacity loss (Cirr=40–100 mAh g−1) or low Coulombic efficiency (usually ≤80%) in the initial charge-discharge cycle due to the structural instability, albeit their relatively high specific capacity [607, 608]. To tackle this issue, various composite materials are made by physically blending these cathode materials with Li-insertion hosts (e.g., V2O5, LiV3O8, Li4Mn5O12 and LiFePO4) to recapture the Li+ ions that cannot be re-intercalated into the cathode structures during the discharge process, instead of solving the underlying causes of the metal ion migrations and oxygen loss [312, 609-612]. 30 wt% Li4Mn5O12 and 18 wt% LiV3O8 have been mechanically blended with the layered Li1.2Mn0.54Ni0.13Co0.13O2 for completely eradicating the irreversible capacity loss, because of the ability of Li4Mn5O12 and LiV3O8 to accept the extracted lithium that could not be inserted back into the layered Li1.2Mn0.54Ni0.13Co0.13O2 [609]. The LiV3O8-

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Li1.2Mn0.54Ni0.13Co0.13O2 composite also showed a high discharge capacity of ∼280 mAh g−1 with little or no irreversible capacity loss and excellent cyclability. A few commercial automotive battery companies have also developed Li-ion batteries with the cathodes composed of a mixture of two or more different cathode materials (e.g., the “Ford Focus” BEV and “Chevrolet Volt” use pouch cells with a blend cathode of NCM and spinel materials), in order to take full advantage of the unique properties of each material and optimize the performance of the battery with respect to the automotive operating requirements [603, 613, 614]. The layered NCA and NCM materials have high capacity and good cycling characteristics, but they are costly and exhibit poor thermal stability. On the contrary, spinel materials show good thermal stability, high voltage, high rate capability and low cost, but they have low capacity. Mixing the two cathode materials could circumvent the shortcomings of the parent materials to achieve higher energy or power density as well as enhanced thermal stability and lower cost [162, 604, 615, 616]. Kitao et al. [617] blended layered LiNi0.4Co0.3Mn0.3O2 with spinel Li1.1Mn1.9O4 (6:4, w/w) (Fig. 29A). Combining NCM with spinel materials resulted in relatively continuous change in the operating voltage region, and the Li-ions were removed from the spinel at high voltage of ≥3.9 V and then from NCM at lower voltages (Fig. 29B) [9]. The mixed electrode also exhibited an average capacity of the spinel and layered electrodes (Fig. 29C). The lattice parameter change of Li1.1Mn1.9O4 in the mixture cathode was different from the original Li1.1Mn1.9O4 cathode during the electrochemical cycling. Moreover, the composite electrode in 18650 full-cells showed better cycling performance than the Li1.1Mn1.9O4 and LiNi0.4Co0.3Mn0.3O2 cathodes after storage for 30 days at 45 oC (Fig. 29D). Myung et al. [618] also found that the mixed positive electrode of spinel Li1.1Mn1.9O4 and LiNi0.80Co0.15Al0.05O2 (1:1 in weight) had a sum average value of the Li1.1Mn1.9O4 and LiNi0.80Co0.15Al0.05O2 electrodes, and it showed greatly improved capacity retention even at 60 oC. The capacity fading was only of ~8 mAh g−1 by applying 1 C rate (150 mA g−1) during 100 cycles. In the full-cell test using graphite as the negative electrode, the retained capacity was ~96% of its initial capacity by applying 1 C current at 25 °C. Aging studies indicated that Mn in Mn-based cathode materials was dissolved from the cathode materials because of the reaction with HF, and the blending of spinel with layered cathode materials (e.g., LiNi0.8Co0.2O2 and NCA) can restrain the Mn dissolution [619, 620].

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Fig. 29 (A) XRD spectra, (B) relationship between the state of charge (SOC) and open circuit voltage (OCV), (C) initial discharge curves, and (D) capacity recovery ratio after storage for 30 days at 45 oC of the single and mixture cathode materials. Adapted with permission [617]. Copyright 2005 The Electrochemical Society.

Olivine-structured cathode materials (e.g., LiFePO4 and LiMnPO4) have better thermal stability than the layered oxides, but they exhibit relatively long and flat chargedischarge voltage plateaus that make them difficult to distinguish the change in state of charge (SOC) for EV applications [604, 613]. Gallagher et al. [604] studied a blend cathode with LiFePO4 and xLi2MnO3·(1–x)LiMO2. Continuous variation of voltage appeared in the cell voltage profile, which is helpful for monitoring the battery SOC and improving the power capability and thermal stability. The addition of LiFePO4 to other electrode materials such as Li1.17Mn0.58Ni0.25O2 and LiCoO2, is also proven to enhance the capacity retention during electrochemical cycling and the rate performance at high discharge currents [621]. When creating the mixed active material cathodes, several parameters must be manipulated such as composition, mass ratio, particle morphology and the physical configuration of the functional materials within the electrode structure [621]. It should be also noted that the interface between the two components affects the charge and mass transfer and is important to the operation of the electrode, so mechanical activation and heat treatment could be utilized to take advantage of the composite properties [622, 623]. NCM has been mixed with LiCoO2 in a composite electrode [624], in which case the

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enhanced performance was ascribed to the effective combination of large and small particles (i.e., the morphology and geometry of the components) rather than the inherent electrochemical properties of the active materials [9]. 4.2. Fabrication process regulation Beside of the conventional slurry coating method, other electrode fabrication methods including vacuum filtering, electrospinning and electro-deposition are also applied for further improving the battery performance [625-628]. Wu et al. [629] developed a free-standing, binder-free NCM/CNT composite film cathode with only LiNi0.5Co0.2Mn0.3O2 particles and 5 wt% single-walled CNTs by vacuum filtering and subsequent drying (Fig. 30A), different from the traditional slurry-coating method with the electrode particles, polymer binders and conductive additives of carbon black (CB). The NCM particles were intimately embedded into the CNT network to form a standalone NCM/CNT composite cathode without the polymer binders and carbon blacks, but random electric contact and crevices of ~300 nm gaping were found in the slurry electrodes of NCM/CNT/CB and NCM/CB (Fig. 30C). Consequently, the binder-free NCM/CNT electrode exhibited lower resistance and higher structural stability than the slurry electrodes (Fig. 30B). The CNT network simultaneously served as an electron transport pathway and also an electrochemically-active cathode ingredient besides LiNi0.5Co0.2Mn0.3O2. Therefore, the polarization of the NCM particles in the binder-free NCM/CNT electrode was significantly reduced, leading to an additional increase in the reversible capacity (262 mAh g−1 vs. 225 and 200 mAh g−1 for the NCM/CNT/CB and NCM/CB slurry electrodes, respectively) and an obvious improvement in the cycling stability and rate performance compared to the slurry-coated electrodes (Fig. 30D-E). Moreover, cyclic voltammograms of these NCM composite cathodes showed that the binder-free NCM/CNT electrode displayed clear well-separated two-stage delithiation processes which cannot be seen for the slurry-coated electrodes, furtherly indicating the depolarization of the CNT-based conductive network (Fig. 30F). In addition to the filtering method, electro-spinning followed by thermal treatment is also usually used to fabricate free-standing and binder-free electrodes [630-632], though the high energy depletion and low production efficiency.

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Fig. 30 (A) Schematic illustration of the fabrication process of free-standing, binder-free NCM/CNT composite film cathodes by vacuum-filtration of the mixture of the NCM particles and CNT in aqueous solution, and (B) resistance between the two sides, (C) SEM images, (D) voltage profiles for the first four charge-discharge cycles between 3.0–4.8 V, (E) rate performance at various current densities of 17–680 mA g−1 and (F) cyclic voltammograms of the as-prepared cathodes. Adapted with permission [629]. Copyright 2014 American Chemical Society.

4.3. Remarks In addition to the materials design on the composition/structure/morphology levels, appropriate cathode design is also sometimes indispensable to improving the cycling performance and thermal stability and lowering the cost for their massive applications. Using proper conductive additives and polymer binders can better disperse the cathode particles, enhance the electrode structure stability, and facilitate the charge transfer. Small amount of novel conductive nano-additives such as carbon nanotubes and graphene with proper size, surface functional groups and electric conductivity can effectively enhance the battery cycling performance. The utilization of conducting binders to replace the conventional PVDF binder can furtherly increase the rate capacity, while the high-thermostability polymer binders can increase the cell safety (or thermal stability) though the partial sacrifice of the cycling performance and high cost. The water-soluble binders such as CMC and alginate show the advantages of low cost, environmental friendliness and simplification of the cell manufacturing; however, the high reactivity between the cathode materials and water causes the dissolution of the metal elements and corrosion of Al current collectors. The combination use of different conductive additives and binders seems a better choice to greatly improve the battery performance. Another approach to adjust the electrode composition is physically blending the layered NCM/NCA particles with other cathode materials (e.g., V2O5,

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Li1.1Mn1.9O4, LiMnPO4 and LiCoO2), and a few companies have applied this method to optimize the overall battery performance (e.g., eliminating the initial irreversible capacity loss, increasing the energy/power densities, and improving the thermal stability and safety) and lowering the whole cathode cost. Some parameters such as composition, mass ratio, particle morphology, materials interface and treatment condition should be carefully controlled for taking full advantage of the properties of each material. Compared to the traditional slurry coating technology, other novel electrode fabrication methods such as filtering and electrospinning are utilized to prepare binder-free, flexible electrodes with improved cycling performance and energy density by increasing the structure stability and electrical conductivity and reducing the inactive mass loading; however, it is still necessary to develop improved or new electrode fabrication methods with high production efficiency and low cost. 5. Electrolyte design The electrolytes in commercial Li-ion batteries consist of the low-flash-point carbonate solvents (i.e., EC/DEC/DMC/EMC), and Li salts (e.g., LiPF6) with high sensitivity to water and temperature. The interactions and reactivity between the active electrode materials, the impurities and the electrolytes would result in the decomposition of the electrolytes (e.g., at ≥4.2 V for the carbonate-based solvents), the dissolution of the metal elements (e.g., Ni and Mn), the corrosion of the Al current collections (e.g., by HF), the exfoliation of the electrode materials, the formation of the thick and unstable SEI films, the growth of Li dendrites, etc. The further use of the NCM/NCA cathode materials with high-content Ni/Co elements, high charge voltage, more Li storage capability and insufficient thermal stability exacerbates most of these issues (Fig. 2 and 4-6), which would seriously affect both the cycling performance and safety of the whole cells. Thus, beside of the NCM/NCA material and electrode designs, the modification of the electrolytes (Fig. 1 and 31) is also becoming more and more important to the applications of the NCM/NCA cathode materials by stabilizing the electrolytes and the electrode materials. The computational study at various molecular levels based on the first-principles calculations offer an efficient and low-cost route to screening electrolytes and understanding the electrolyte properties, but the experiments are indispensable to the validation of the computational results because of the complicated real situations and the limited computational capabilities of modern computers [633, 634].

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Fig. 31 Electrolyte design methods.

5.1. Li salts The electrolyte conductivity is closely related to the variety of Li salts and the solvents and their proportion (Table 3). There are a few important inorganic Li salts such as LiClO4, LiAsF6, LiBF4 and LiPF6. LiPF6 is always used in the commercial electrolytes, because of its high ionic conductivity, low oxidizability and low cost. The formation of protective aluminum (oxy-)fluoride layers on the Al current collectors can prohibit the anodic Al corrosion. Nevertheless, it can easily decompose with trace water to generate LiF, PF5, OPF3 and HF, which would facilitate the metal dissolution from the NCM/NCA cathode materials and exacerbate unwanted side reactions. The LiPF6resulted SEI films also have high resistance at low temperature, and low stability at high temperature (easily decompose at ≥80 oC). Thus, the modification of LiPF6 and development of novel Li salts are becoming urgent. A few new groups are usually used to replace the F atoms in LiPF6 (e.g., LiPFm(C2F5)6‒m), in order to improve the thermal stability and ionic conductivity of LiPF6-resulted SEI films. Lithium bis(fluorosulfonyl)imide (LiFSI) and LiTFSI have high conductivity and thermal stability, but their large anion size results in the high electrolyte viscosity. Additionally, the corrosion of Al current collectors by the generation of Al(FSI)x and Al(TFSI)x impedes their applications in organic liquid electrolytes. Lithium bis(oxalato)borate (LiBOB) has high ionic conductivity, wide electrochemical window and high thermal stability, but it is sensitive to water. Lithium difluoro(oxalato)borate (LiDFOB) not only possesses high conductivity, thermal stability and good compatibility with the graphite anodes (by forming dense SEI films at ~1.5 V), but also can suppress the oxidation of the electrolytes and the dissolution of the

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metal elements. Park et al. [635] found that the Al corrosion inhibition ability of Li salts in EC/DEC (3/7, v/v) was in the order of LiDFOB > LiBF4 ≈ LiPF6 > LiBOB, and the superior inhibition ability of LiDFOB was ascribed to the formation of a passive layer of Al‒F, Al2O3 and B‒O species on the Al surface, which made the full cells with 0.8 M LiFSI and 0.2 M LiDFOB exhibit comparable rate capacity to that of the 1 M LiPF6based cells. Nevertheless, LiDFOB has no competitive advantage with respect to price. Table 3 The commonly-used Li salts in Li-ion batteries Type

Decomposition temperature (oC) 125

Conductivity (mS/cm) 10.8 (EC/DMC, 1:1 v/v)

Pros

Cons

LiPF6

Molecular weight 152

High conductivity, excellent roomtemperature performance, low cost

LiFAP

452

200

Without decomposition with water, high-temperature stability

LiFSI

187

200

8.2 (EC/DMC, 1:1 v/v) 9.7 (EC/DMC, 3:7 v/v)

Decomposition with water, poor performance at high or low temperature High cost

LiBETI

387

250

LiTFSI

287

360

LiBOB

194

275

LiDFOB

144

200

11.1 (EME/DOL, 1:1 v/v) 9.0 (EC/EMC, 1:1 v/v) 14.9 (DME)

8.6 (EC/DMC, 1:1 v/v)

High conductivity, wide electrochemical window, high dissolution ability, indecomposition with water, excellent performance at high or low temperature High conductivity, high thermal stability, high stability of the SEI films High conductivity, high dissolution ability, high thermal stability

Corrosion of Al foils at high voltage

High conductivity, wide electrochemical window, high thermal stability, high stability of the SEI films High dissolution ability, wide electrochemical window, high performance at high or low temperature, high-stability SEI films

Poor dissolution ability, sensitive to water

High cost

Corrosion of Al foils at high voltage

High cost

5.2. Solvents In theory, ideal electrolyte solvents should have high dielectric permittivity and low viscosity, which can make Li salts have high dissolution ability and dissociation degree in the solvents. However, the cyclic alkyl carbonates such as ethylene carbonate (EC) and propylene carbonate (PC) with high permittivity have high viscosity due to the interaction between the molecules, while the linear alkyl carbonates such as dimethyl carbonate (DMC), diethyl carbonate (DEC) and ethyl methyl carbonate (EMC) with low viscosity show low permittivity (Table 4). So the mixture solutions of high-permittivity cyclic carbonates and low-viscosity linear carbonates with certain proportions (e.g., EC/DEC and EC/DMC with a volume ratio of 1:1) are usually used as electrolyte solvents. Unlike PC, EC cannot get into the graphite anodes but form effective SEI

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films (~0.8 V) to inhibit the further solvent decomposition and the insertion of solvents into the electrode materials. Density functional theory (DFT) calculations and molecular dynamics (MD) simulations also prove that the SEI films have higher density, more cohesive energy, less solubility in solvents, and less favorable interactions with graphite surface than the films from the PC decomposition [636, 637]. Thus, EC is an indispensable solvent in commercial electrolytes. Table 4 The commonly-used organic solvents in Li-ion batteries Type Propylene carbonate (PC) Ethylene carbonate (EC) Diethyl carbonate (DEC) Dimethyl carbonate (DMC) Ethyl methyl carbonate

Structure

Melting point (oC)

Boiling point (oC)

Permittivity (F/m, 25 oC)

Viscosity (mPa·S, 25 oC)

Flash point (oC)

Density (g/mL, 25 oC)

–48.8

242

64.92

2.53

132

1.200

36.4

248

89.78

1.90 (40 oC)

160

1.321

–43

126

2.81

0.75

31

0.969

4.6

91

3.11

0.59 (20 oC)

18

1.063

–53

110

2.96

0.65

26.7

1.006

（EMC）

The electrochemical oxidative stability of the solvent molecules was studied by Xing et al. based on DFT calculations [638]. The highest occupied molecular orbital (HOMO) energy of EC, PC, DMC, DEC and EMC is –8.46, –8.37, –8.21, –8.05 and – 8.11 eV, respectively, indicating that the oxidative decomposition activity is in the order of DEC > EMC > DMC > PC > EC. Nevertheless, PF6– in LiPF6 is more easily coordinated with cyclic alkyl carbonates (e.g., EC) than linear alkyl carbonates, because of the higher dielectric constant of cyclic alkyl carbonates than linear alkyl carbonates. The binding energy of EC-PF6– (–64.51 kJ/mol) is more negative than others, further proving that EC-PF6– would reach cathode and oxidized more easily in the battery charge process. They also found that PF6– and ClO4– anions in Li salts can significantly lower the oxidation stability of EC and affect the oxidation decomposition paths [639]. Again one should note that the theoretical calculations of oxidation/reduction reactions of single molecule may not reflect the real situations owing to the mixture of salts and solvents polarized at the electrode interfaces and their complicated interactions, and the computational results should also be validated by the experiments [640, 641]. To widen the electrochemical window of the electrolytes for high-voltage

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NCM/NCA batteries, researchers are exploiting high-concentrated electrolytes such as the mixture solution of propylene carbonate (PC), γ-butyrolactone (GBL) and dimethylcarbonate (DMC) where most of the solvent molecules coordinate Li+ ions in the high-concentration electrolytes. Cao et al. [642] found that the LiNi0.8Co0.1Mn0.1O2/Li half-cells with nearly saturated 8.67 mol/kg LiBF4/DMC electrolytes exhibited higher capacity retention of 93.3% than the conventional electrolyte-based cells (i.e., 1 M LiPF6:EC/DMC) after 50 cycles at 0.1 C within 3.0‒4.3 V, because of the effect suppression of the electrolyte decomposition and cathode particle pulverization. Novel electrolyte solvents with higher oxidation potentials than EC are also requested to elevate the charge potential of the NCM/NCA cathode materials for highdensity Li-ion batteries. Ionic liquids (ILs) have wide electrochemical window, low volatility and high flame retardancy, and are thus considered as high-voltage and safer solvent candidates. ILs are usually comprised of quaternary ammonium cations such as pyrrolidinium (PYR), piperidinium (PP), pyridinium, imidazolium and ammonium derivatives, and inorganic or organic counteranions such as PF6‒, BF4‒, [(FSO2)(CF3SO2)N]‒, [(CF3SO2)2N]‒, [(C2F5SO2)2N] ‒ and [(CF3SO2)(CF3CO)N]‒ [643]. However, their high viscosity, low ionic conductivity and low wettability lead to the poor cycling performance of the IL-based cells. Most IL-based electrolytes also show compatibility problems with the graphite anodes [644]. A few carbonate solvents or functional additives are commonly mixed with the ILs to facilitate the practical applications of these ILs in electrolytes [645]. Fluorinated organic solvents with high-electronegativity and low-polarity fluorine atoms such as 3-(1,1,2,2-tetrafluoroethoxy)-1,1,2,2-tetrafluopropane (FEPE), 1H,1H,5H-perfluoropentyl-1,1,2,2-tetrafluoroethylether (FEAE), methyl(2,2,2trifluoroethyl)carbonate (FEMC) and di-(2,2,2 trifluoroethyl) carbonate (DFDEC) are utilized as solvents for high-voltage electrolytes [646]. Other organic solvents including nitriles and sulfones were also exploited. However, the application of these new solvents are restricted by the high viscosity, poor wettability of the olefin separators, high gas production, low conductivity, high cost or safety issue [647]. Elevating the oxidation potential with the fluorine solvents usually raise the reduction potentials and is detrimental to graphite anode side or other battery components, though the improvement of the thermal stability and high-voltage cycling performance of the NCM811/Li-based half-cells with the electrolytes such as LiPF6:PC/FEMC/DFDEC [648]. Improving the bulk properties of the electrolytes by the single use of novel solvents is probably not an appropriate approach. Zhang et al. [649] also reported highperformance NCM523/Li half-cells using 1 M LiBF4:FEC/succinonitrile (SN) to replace 1 M LiPF6:EC/DEC/DMC electrolyte. The FEC/SN can form thinner, more uniform, Ncontaining high-conductivity protective layers on the NCM particles and prevent the

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electrolyte decomposition and harmful side reactions, and thereby the FEC/SN-based cells showed higher capacity retention of 73.6% than the EC/DEC/DMC-based cells (35.5%) after 100 cycles within 3.0–4.7 V. Removing EC from the typical organic carbonate-based electrolytes by adding small amounts of electrolyte additives is proven to be a feasible way to obtaining better battery performance than those based on the EC-based electrolytes [650, 651]. Dahn group [647] found that the removal of EC in the electrolyte (EMC as solvent with a few additives such as VC and FEC) can enhance high-voltage performance of NCM424/graphite cells at both room temperature and high temperature (e.g., the cells using 98% EMC and 2% FEC showed an increased energy density by at least 10% with a cutoff voltage of 4.4 V), and adjusting the loading of the additives can also effectively passivate the graphite anode, ensure low impedance and gassing, lower the polarization during the high-voltage cycling, and improve the safety. He et al. [652] compared the effect of a novel 1.0 M LiPF6:FEC/HFDEC/1% LiDFOB electrolyte with the conventional 1.2 M LiPF6:EC/EMC electrolyte, and found that the novel electrolyte showed a higher oxidation peak of 7.05 V than the conventional electrolyte. The formation of SEI films on the NCM523 cathode and graphite anode by the decomposition of LiDFOB and FEC would protect the electrodes from further electrolyte corrosion, respectively, and thereby the novel electrolyte-based NCM523/graphite cells displayed a higher capacity retention of 82% than the conventional electrolyte-based cells (67%) after 100 cycles at C/3 within 3.0‒4.6 V. 5.3. Additives Beside of the modification on Li salts and solvents, functional electrolytes are exploited to protect the electrolytes and the electrode materials. The concept of “functional electrolytes”, i.e., high-purity electrolytes with a slight amount (≤10 wt%) of electrolyte additives was proposed by Yoshitake et al. in 1999 [653]. A few electrolyte additives such as propargyl methanesulfonate (‒0.22 eV) and propargyl methyl carbonate (0.82 eV) have lower LUMO energy than EC (1.18 eV) and DMC (1.05 eV) and can be reduced at higher potentials at 1.24 V and 0.83 V, respectively [653], and then form SEI films with high Li conductivity and stability against electrolyte attack, which would prevent the exfoliation of the graphite electrodes and improve the battery performance [654]. Other additives such as pyridine-boron trifluoride (PBF) and prop-1-ene-1,3-sultone (PES) can effectively inhibit the impedance growth and also the gas production during the battery formation and cycling [655]. According to the functionality, the electrolyte additives are classified as SEI film formation additives, acid scavengers, flame retardants, overcharge protectors, wettability additives, etc. Although the single use of electrolyte additives cannot tackle the root problem of the NCM/NCM ternary materials, the exploitation of appropriate functional additives is an effective and economic method for improving the cycling

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performance and battery safety. 5.3.1. SEI film formation agents Similar to the functional additives for anodes such as vinylene carbonate (VC), fluoroethylene carbonate (FEC) [656], difluoroethylene carbonate (DiFEC) [657], 2,3epoxypropyl methanesulfonate (OMS) [658] and dicyano ketene vinyl ethylene acetal (DCKVEA) [659], a few other functional additives such as methyl(2,2,2-trifluoroethyl) carbonate (FEMC), isophorone diisocyanate (IPDI) [660], ethylene sulfite (ES), divinyl sulfone (DVS) [661], phenyl vinyl sulfone (PVS) [662], 3,3’-sulfonyldipropionitrile (SDPN) [663], triallyl phosphite (TAPi) [664], trimethylboroxine (TMB) [397], triethylborate (TEB) [665], di(methylsulfonyl) methane (DMSM) [666], tris(trimethylsilyl) borate (TMSB), [4,4’-bi(1,3,2-dioxathiolane)] 2,2’-dioxide (BDTD) [667], tris(trimethylsilyl) phosphate (TMSPa) [668, 669], tris(2,2,2-trifluoroethyl) phosphite (TTFP) [670], allylboronic acid pinacol ester (ABAPE) [671], and 3,3'(ethylenedioxy)dipropiononitrile (EDPN) [672] with lower oxidation potentials than the base electrolytes (e.g., EC/DEC/EMC) [673] can also decompose before the baseline electrolytes and form stable SEI films to protect the NCM/NCA cathode materials from further destruction and suppress the electrolyte decomposition [674-679]. It should be noted that the resultant films from the decomposition of the organic additives are mainly composed of low-conductivity polymers, which may increase the interface resistance between the electrodes and the electrolytes. Zheng et al. [680] found that the ionic conductivity of the electrolyte (1 M LiPF6:EC/DEC/EMC) decreased from 11.8 to 10.2 mS/cm with the addition of 2 wt% N-allyl-N,N-bis-(trimethylsilyl)amine (NNB), but the generation of SEI films from the NNB decomposition (at ~3.9 V in the LSV curve) can effectively inhibit the subsequent electrolyte oxidation and NCA destruction, which made the capacity retention increase from 72.8% to 86.2% after 300 cycles at 1 C within 3.0‒4.2 V. Liao et al. [681] investigated the effect of TMSB on the selfdischarge behavior of the charged LiNi1/3Co1/3Mn1/3O2 under a high potential of 4.5 V by DFT calculation and experiment measurements. TMSB had a higher HOMO energy but a lower oxidation potential (5.76 V) than EC (6.96 V) and DMC (6.84 V). The preferential electrochemical oxidation of TMSB in the EC/DMC electrolyte during the cycling can effectively inhibit the decomposition of the electrolyte and protect the chemical reactivity between the electrolyte and NCM333, and thus suppress the selfdischarge of the charged NCM333 cathode. Qian et al. [682] scrutinized the effect of VC, ES and FEC on the SEI thickness on the NCM333 cathode during the different electrochemical periods: after formation (0.1 C for 5 cycles) and long-term cycling (0.1 C for 5 cycles, and then 1 C for 100/150 cycles). Except the 2 vol% FEC-based cells, most of the electrolyte decomposition productions generated on the cathodes during the aging period. The 2 vol% FEC-based electrode showed the thinnest SEI thickness of 0.8 nm after the long-term aging, but the SEI thickness on the 2 vol% VC-based cathode

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increased from 1.0 to 2.9 nm when cycled from 100 to 150 cycles. The 2 vol% FECbased cells also exhibited the highest capacity of 114 mAh g−1 after 150 cycles. In very contrast to the abovementioned additives, a few organic molecules such as dopamine and 2-(2,2,2-trifluoroethoxy)-1,3,2-dioxaphospholane 2-oxide (TFEOP) [683] can be easily electrochemically oxidized and polymerized at lower potentials than the electrolytes to form the polymer-based passive films or SEI films on the cathode surfaces and therefore prevent the electrolyte decomposition and remain the electrode stability. Lee et al. [684] calculated the HOMO and LUMO energy of the dopamine additive and the solvents of EC, EMC and DMC, and found that dopamine had lower LUNO energy (‒0.96 eV) and higher HOMO energy (‒11.06 eV) than the solvents (around 1.5 and ‒13 eV). This result suggested that dopamine should decompose before the electrolyte to form a passive polydopamine film on the cathode, which was proved by LSV and XPS measurements. When the NCM333/graphite cells were chargedischarged between 3.0‒4.5 V at 1 C, the dopamine-contained cells exhibited higher capacity of 147 mAh g−1 and capacity retention of 90.1% than the dopamine-free cells (136 mAh g−1 and 83.3%) after 100 cycles. EIS results showed that the dopaminecontained cells had much lower passivation layer resistance and charge transfer resistance than the dopamine-free cells, and XPS measurements of the cycled cathodes showed that the dopamine-contained cathode had more intense C-F peak and weaker LiF peak, proving the effect inhibition of the electrolyte (solvents and PF6‒) decomposition by the passive polydopamine layer. In addition to the organic electrolyte additives, a wide variety of Li-based salts including LiBOB, LiDFOB, lithium difluorophosphate (LiDFP) [685, 686] and lithiumcyclo-difluoromethane-1,1-bis(sulfonyl)imide (LiDMSI) [687] have also been used as SEI film forming additives at the NCM/NCA cathodes. Hong et al. [688] compared the effects of prop-1-ene-1,3-sultone (PES) and LiDFP on the NCM523/graphite cells with 1 M LiPF6:EC/EMC electrolyte, and found that the 1 wt% LiDFP-containing cells displayed higher capacity of 177.0 mAh g−1 and capacity retention of 90.2% than the 1 wt% PES-containing (161.0 mAh g−1 and 77.1%) and additive-free (113.1 mAh g−1 and 64.0%) cells after 200 cycles at 1 C within 3.0‒4.45 V, due to the formation of thinner and high-conductivity interface layers on the cathode surfaces by the LiDFP decomposition. Kaneko et al. [689] studied the composition of the LiBOB and LiDFOB-resulted SEI films extracted from the cathode material surface by XPS, ICP and NMR, and found that the SEI films were composed of oxalate groups and bis(4,5dioxo-1,3,2-dioxaborolan-2-yl) oxalate (BDBO). The quantum mechanics (QM) and molecular mechanics (MM) methods were further applied to disclose the redox nature and oxidative decomposition dynamics. The higher average HOMO energies of the anions of LiBOB (‒9.2 eV) and LiDFOB (‒10.2 eV) than EC (around ‒11 eV) indicated the first oxidative decomposition of LiBOB and LiDFOB on the cathode surface (Fig.

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32A-B). MM further proved the generation of CO2 gas from the oxalate groups with the B-O bond breaking. Moreover, the free energy profiles showed that the BDBO-induced SEI phases had high stability and Li+ ion translocation ability (Fig. 32C-D), which resulted in the improvement of the cathode performance.

Fig. 32 Energy distributions of HOMO/LUMO in (A) LiBOB- and (B) LiDFOB-based electrolytes (EC as solvent). (C) Equilibrium snapshot of the SEI/electrolyte interface. Purple transparent shell and yellow ball represent SEI phase and Li ion, respectively. All of the chemical compounds are in stick form. (D) Free energy profiles along the interface. Red and green lines represent the free energies of BDBO and Li ion detaching from the interface, respectively. Adapted with permission [689]. Copyright 2016 Elsevier.

A few LiPF6 hydrolysis products such as dimethyl fluorophosphate (DMFP) and diethyl fluorophosphate (DEFP) generate during the battery cycling especially at high operation voltages or temperatures, and can also act as effective SEI film formation additives though their high toxicity. Wagner et al. [690] found that DMFP and DEFP with Li attaching to F or O had lower oxidative stability of ~2.5 eV than dimethyl methyl phosphonate (DMMP, 5.1 eV) electrolyte, and further proven the easy oxidative decomposition of DMFP and DEFP to form SEI protection layers on the cathodes by CV and XPS measurements. As a result, the NCM333/Li cell with the DMFP-involved LiPF6:EC/EMC electrolyte displayed higher discharge capacity of 159.0 mAh g−1 and capacity retention of 92.7% than the additive-free cells (140.3 mAh g−1 and 83.8%) after 50 cycles at 30 mA g−1within 3.0‒4.6 V. Apart from the aforesaid additives which can only be electrochemically oxidized on the cathode side or reduced on the anode side, researchers are also developing novel

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additives such as ethyl 4,4,4-trifluorobutyrate (ETFB) [691], lithiumbis(hexafluorobutane-2,3-diol)-borate (LiHFBDB) [692], lithium 4,5-dicyano-2(trifluoromethyl)imidazolide (LiTDI) [693], and lithium fluoromalonato(difluoro)borate (LiFMDFB) [694] that can simultaneously form cathode and anode interface films with low impedance before the bulk electrolyte components during the battery cycling process. Liao et al. [695] first proved the easy oxidation and reduction of LiDFBOP at lower potentials than the electrolyte solvents (EC/EMC) by HOMO/LUMO energy calculations (Fig. 33A) and LSV measurements (Fig. 33B-C). The CV measurements also proved the high ion conductivity and fast Li insertion/extraction kinetics of the (Fig. 33D) cathode and (Fig. 33E) anode interface films derived from the oxidation and reduction of LiDFBOP, respectively. The additive-derived films can effectively inhibit the electrolyte decomposition, protect the electrodes from destruction and facilitate the Li transfer. Therefore, the LiDFBOP-containing LiNi0.5Co0.2Mn0.3O2/graphite pouch cells exhibited higher capacity retentions (75% and 93%) and Coulombic efficiencies than the additive-free cells (21% and 73%) at both (Fig. 33F) 25 and (Fig. 33G) 0 oC. TEM also proved the formation of thinner SEI films on the LiDFBOP-containing NCM particles after the low-temperature cycling (Fig. 33H-I).

Fig. 33 (A) The calculated HOMO/LUMO energy of solvents and electrolyte additives, LSV curves of platinum electrodes in electrolytic cells using baseline (1 M LiPF6:EC/EMC) and 1% LiDFBOPcontaining electrolytes on the (B) cathode and (C) anode sides, mechanisms of the formation of (D)

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cathode and (E) anode interface films from LiDFBOP, cycling performance of LiNi0.5Co0.2Mn0.3O2/graphite pouch cells at (F) 25 and (G) 0 oC, and TEM images of the NCM particles (H) without and (I) with LiDFBOP additive after the 0 oC cycling. Adapted with permission [695]. Copyright 2018 Wiley-VCH.

5.3.2. Acid scavengers The generated acid (e.g., HF) in the LiPF6-based electrolytes with the presence of H2O would destroy the SEI films and dissolve the Ni/Co/Mn elements in the NCM/NCA cathode materials. Thus, a few inorganic metal oxides and carbonates such as Al2O3, MgO, BaO and CaCO3 are used as acid scavengers to remove HF by the reaction with HF, maintaining the structure stability of the cathode materials. A few organic compounds such as tris(trimethylsilyl) borate (TMSB) [696], tris(trimethylsilyl) phosphite (TMSPi), tris(trimethylsilyl) phosphate (TMSPa), hexamethylphosphoramide (HMPA), tris-2,2,2-trifluoroethyl phosphite (TTFP), dimethoxydimethylsilane (DODSi) [697, 698], diethyl phenylphosphonite (DEPP) [699], N,N-diethylamino trimethylsilane (DEATMS) [700] and p-toluenesulfonyl isocyanate (PTSI) [701] can also react directly with PF5, HF, F‒ and PF6‒ in the electrolytes, inhibiting the formation of LiF and Li2CO3, and the dissolution of metal elements. Han et al. investigated the role of TMSPi with the combination of DFT calculation and experiments [702]. TMSPi had higher HOMO and LUMO energies of ‒6.52 and 0.70 V than EC (‒8.25 and 0.60 V), respectively, and thus exhibited lower oxidation and reduction potentials of 4.29 and ‒1.03 V than EC (6.92 and ‒0.32 V). TMSPi also had a higher reactivity with HF (10.4 kcal/mol) than TMSPi+ (3.2 kcal/mol), and thus may react with HF rather than as a cation on the cathode surface. It was proven that TMSPi can remove HF in the electrolyte with the reaction with HF and form a protection film on the cathode surface to improve the battery performance. They furtherly investigated the self-decomposition reactions and reactivity of TMSB, TMSPi and TMSPa with HF and LiF through first-principles calculations [703], and found that these additives cannot spontaneously decompose even in the cation forms but can effectively react and scavenge HF and LiF in the electrolyte and on the cathode surface, thus improving the cathode/electrolyte interface stability and battery performance. Deng et al. [704] found that diphenyldimethoxysilane (DPDMS) with Si-F and Si-O bonds can react with trace amounts of HF and PF5 in the electrolyte to form a surface film on the cathode and protect the cathode from electrolyte corrosion, and therefore the LiPF6:EC/EMC/DMC/1 wt% DPDMS-containing NCM622/Li cells showed a higher capacity retention of 93.3% than the DPDMS-free cells (71.9%) after 200 cycles at 2 C and 55 oC within 2.8‒4.3 V. Other organic additives can directly react with H2O in the electrolyte to suppress the HF generation and meanwhile form SEI layers on the cathode surfaces, and thereby

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maintain the structural stability of the cathode materials. Liao et al. [705] studied the effect of 1-(2-cyanoethyl) pyrrole (CEP) additive on the NCM622 cathode in 1 M LiFP6:EC/EMC. CEP showed a higher HOMO energy of ‒6.15 eV than the solvents of EC (‒8.73 eV) and EMC (‒8.40 eV, Fig. 34A), and thus the 1 wt% CEP-containing electrolyte exhibited a lower oxidation potential than the baseline electrolyte (Fig. 34B). The resulted SEI films from the oxidation of CEP in the initial charge process (Fig. 34C) can effectively prohibit the decomposition of the electrolyte. CEP also had higher combination energy with HF (‒32.47 kJ/mol) and H2O (‒51.87 kJ/mol) than EC and EMC (Fig. 34D-E), and would prevent HF generation by eliminate H2O (Fig. 34F) and the transition metal dissolution by HF corrosion (Fig. 34G). Therefore, the 1 wt% CEPcontaining NCM622/graphite cells showed higher discharge capacity and capacity retention (81.5%) than the CEP-free cells (27.4%) after 50 cycles at 1 C within 3.0‒4.35 V (Fig. 34H).

Fig. 34 (A) Calculated HOMO energy levels of CEP, EMC and EC, (B) LSV curves of the platinum electrode in the baseline (1 M LiPF6:EC/EMC) and 1% CEP-containing electrolytes, (C) oxidation mechanism of CEP, combination energy of CEP, EC and EMC with (D) HF and (E) H+/H2O, (F) illustration on the coordination of CEP with H2O, (G) contents of Ni/Co/Mn deposited on the cycled graphite anodes, and (H) cycling performance of NCM622/graphite cells at 1 C within 3.0‒4.35 V. Adapted with permission [705]. Copyright 2018 American Chemical Society.

5.3.3. Flame retardants Conventional organic solvents in the commercial electrolytes are also combustion fuels particularly when the batteries are under high current, overcharge, overheating or damaged, and would cause thermal runaway, fire or explosion [706-708]. With the applications of high-energy-density cells or high-capacity battery packs for especially

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EVs, the safety issues are motivating researchers the development of non-flammable organic solvents, flame-retardant electrolyte additives, and even all-solid-state batteries. Nevertheless, the energy density and cycling stability should not compromise the safety issues. As we have mentioned in Section 5.2, the use of fluorinated solvents can elevate the cell potential, but would cause other problems and finally have negative impact on the battery performance. Likewise, the utilization of nonflammable solvents such as ionic liquids, organosilicon compounds, hydrofluoroethers (HFEs) and phosphates [709] can also enhance the battery safety, but rational design of the whole electrolyte is needed to maintain the pristine battery performance. Pham et al. [710] developed a nonflammable electrolyte of 1 M LiPF6:PC/di-(2,2,2 trifluoroethyl)carbonate (DFDEC)/FEC and studied its impact on the flame resistance and cycling performance of Li1.13Ni0.203Co0.203Mn0.463O2/graphite cells. The co-intercalation of PC with Li+ ions into the graphite interlayers can be suppressed by the use of 1 wt% FEC additive, which can form protective SEI layers on the graphite surface during the battery charge process. The conventional EC/EMC electrolyte would easily catch fire and show a longer selfextinguishing time of 60 s/g than the PC/DFDEC/FEC electrolyte (<6 s/g). Because of the formation of stable SEI films on the cathode and anode surfaces by the decomposition of DFDEC and FEC and the protection of the electrode materials from electrolyte corrosion, the PC/DFDEC/FEC-containing NCM/graphite cells showed a higher capacity retention of 80% than the PC/DFDEC- (66%) and EC/EMC-containing cells after 100 cycles at 0.2 C within 2.5‒4.85 V. The PC/DFDEC/FEC-containing cells with the charged potential of 4.85 V also exhibited higher voltage retention of 89% than the PC/DFDEC- (66%) and EC/EMC-containing (23%) cells after 14 days at 60 oC. Apart from the use of the nonflammable solvents, a few fire-retardant additives are also developed to eliminate the safety hazards of the NCM/NCA-based Li-ion batteries. The functional additives can generate radical species (e.g., PO2· or HPO2·) at elevated temperature, and then effectively capture the free radical species (e.g., H· or ·OH) that are formed due to the thermally-induced decomposition of the organic carbonate solvents and would cause cascading-chain-propagation gas-phase combustion reactions [711]. The alkyl phosphorus-based compounds such as trimethylphosphate (TMP), triethylphosphate (TEP), dimethyl methyl phosphonate (DMMP), hydrophobic triphenyl phosphate (TPP) and tributylphophate (TBP) can release phosphoric acid and form char and thick glassy carbon layer, which can capture the generated free radicals during the electrolyte combustion process and suppress the liberation of flammable gas [712]. Nevertheless, most of the phosphorous-based flame retardants had drawbacks such as high viscosity and environmental pollution. The instability at low potentials and ineffectiveness to form stable SEI films on the electrodes also hinder their practical applications [643, 709, 713]. The fluorinated phosphates such as tris-(2,2,2-trifluoethyl)

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phosphate (TFP), bis(2,2,2-trifluoroethyl) methyl phosphate (BMP), 2,2,2-trifluoroethyl diethyl phosphate (TDP), pentafluoro(phenoxy)cyclotriphosphazene (FPPN) and ethoxy (pentafluoro) cyclotriphosphazene (PFPN) can not only guarantee the high conductivity of the electrolytes and excellent cycling performance of the batteries by forming stable SEI films on the graphite anodes [714], but also exhibit higher flame resistance than the alkyl phosphates [715, 716]; however, the halogen-based flame retardants are harmful to the environment and human bodies [717, 718]. 5.3.4. Overcharge inhibitors The strict control of the voltage runaway is becoming more and more important to the battery safety, because overcharge leads to unwanted high-voltage and the serious consequences such as short-circuit, gas generation, thermal runaway, explosion and fire [719]. Apart from the battery management system (BMS) and the cell construction design such as pressure valve and current breaker, overcharge protection additives also play an important role in inhibiting the overcharge of the high-voltage NCM/NCAbased batteries, preventing the structural degradation of the cathode materials and electrolyte decomposition, and improving the battery safety. This can also simplify the cell manufacture process and lower the cost. According to the overcharge protection mechanisms, the overcharge inhibitors are classified as the shutdown-type electric polymer additives and redox shuttle additives [720]. The shutdown-type additives usually consist of biphenyls, cyclohexyl benzenes, esters, N-phenyls and their derivatives. These additives can polymerize at certain potentials and form insulating polymer films on the cathodes, which can effectively prohibit the further increase of the cell potential during the charge process. Nevertheless, the general irreversibility of the shutdown process and the potential increase of the internal pressure in the cell related to the polymer layer formation are the main drawbacks of these additives [643]. The redox shuttle additives usually consist of metallocene compounds, polypyridine complexes, thiophene anthracenes, fennel benzenes, anisoles and their derivatives. These additive molecules (RS) can be reversibly oxidized at defined potentials on the cathode side (RS → RS+ + e‒), and thus prohibit the overcharge of the cells by working as overcharge current shunts. The oxidized additives can also easily diffuse to the anode side, and then be reversibly reduced to their initial states (RS‒ ‒ e‒ → RS) [721]. To meet the practical applications of these additives, a few requirements must be fulfilled: (1) the oxidation potential of the additives should be slightly higher than the cutoff charge potential of the cathodes (usually 0.1‒0.2 V); (2) the redox molecules and their oxidization/reduction products should be thermally stable and chemically inert against all the cell components; (3) the additives should have high solubility and diffusion coefficients in the electrolytes; (4) the oxidation-diffusion-reduction-diffusion cycle should be highly reversible [643]. However, many of these redox shuttle additives have

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negative impact on the cell performance under the normal operation conditions [722]. Janssen et al. [722] investigated the effects of 1,3-dimethylimidazolidin-2-mmtrifluoroborate (NHC-BF3), 1,3-dimethylimidazolidin-2-mmtetrafluorotrifluoromethylphosphate (NHC-PF4CF3) and 1,3-dimethylimidazolidin-2mm-pentafluorophosphate (NHC-PF5) on the overcharge behavior and cycling performance of NCM333/Li cells. NHC-BF3 and NHC-PF4CF3 were identified as promising overcharge additives for NCM333 electrodes up to 4.4 and 4.6 V vs. Li/Li+, respectively, and these additives showed no big negative influence on the cell performance under normal charge/discharge conditions within 3.0‒4.2 V (Fig. 35A-B). The shutdown effect was attributed to the decomposition of the additives on the cathodes rather than in the electrolyte/separator phase, but the thin decomposition layers on the cathodes would hinder the intercalation/de-intercalation of Li+ ions into the NCM333 cathodes (Fig. 35C-D).

Fig. 35 (A) Cycling performance and (B) potential profiles of NCM333/Li cells with baseline and 0.1 M NHC-BF3 additive-based electrolytes, (C) AFM images of the cycled NCM cathodes with the NHC-BF3 additive-based electrolyte, and (D) cycling performance of the harvested NCM/Li cells after overcharge using 0.1 M NHC-BF3 additive in 1 M LiPF6:EC/DEC electrolyte. Adapted with permission [722]. Copyright 2017 Elsevier.

5.3.5. Multifunctional additives As mentioned above, most of the additives exhibit single specific function, and thus two or several different additives are needed to improve the overall properties of the electrolytes/cells; however, the increased content and complicated interactions between the functional additives and the cell components may have negative impacts on the cell performance and safety [723]. So a few bi-functional and multifunctional additives are furtherly developed to enhance the overall electrolyte/cell performance. For example, 4bromo-2-fluoromethoxybenzene (BFMB) was used as a novel bi-functional additive,

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which can electrochemically polymerize at 4.6 V (vs. Li/Li+) to form a polymer film for preventing the voltage rise when overcharging and also can lower the flammability of the electrolyte and enhance the thermal stability of the batteries. Zhu et al. [724] designed a single molecule additive of diethyl(thiophen-2-ylmethyl)phosphonate (DTYP) composed of thiophene and phosphate groups with the combined functions of SEI formation agents, acid scavengers and flame retardants: (1) the thiophene groups can polymerize to form ionically conductive cathode interface films; (2) the oxygen of the phosphate groups can capture PF5 in the electrolyte through the acid-base coordination interaction; (3) the phosphite groups can terminate the radical propagation in the combustion process (Fig. 36A). Because of the formation of stable and conductive SEI films on the LiNi0.5Mn1.5O4 cathodes and effective inhibition of HF corrosion, the 0.5 wt% DTYP-containing cells showed higher capacity retention of 84% than the 0.5 wt% thiophene (TH)-containing (79%) or additive-free (18%, 1 M LiPF6:EC/EMC) cells after 280 cycles at 1 C and 60 oC within 3.0‒4.9 V (Fig. 36B). The calculated relative energy of PF5–DTYP (–7.55 kJ/mol) and PF5–P2O5 (–38.13 kJ/mol) were higher than PF5–EMC (–0.372 kJ/mol) and PF5–EC (–0.428 kJ/mol, Fig. 36C), and the 0.5 wt% DTYP-containing electrolyte also showed higher thermal decomposition onset temperature of 223 oC than the baseline electrolyte (193 oC), suggesting the high thermal stability of the DTYP-based electrolyte and the contribution from the generated P2O5 from the DTYP decomposition. Moreover, the Celgard 2400 separators wetted with the DTYP-containing electrolyte cannot be ignited, while the separators itself or soaked in the base electrolyte were easily flammable, indicating the high flame resistance of the DTYP additive (Fig. 36D).

Fig. 36 (A) Molecule structure and functionalities of DTYP additive, (B) cycling performance of the

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LiNi0.5Mn1.5O4 electrodes at 1 C and 60 oC within 3.0‒4.9 V, (C) relative energies and optimized structures of PF5–EC, EMC, DTYP and P2O5, and (D) flammability of the Celgard 2400 separators with baseline and DTYP-containing electrolytes. Adapted with permission [724]. Copyright 2018 The Royal Society of Chemistry.

5.4. Remarks Apart from the designs on the NCM/NCA materials and electrodes, the further modification on the electrolytes offers another important way to stabilizing the electrode/electrolyte interfaces and enhancing the thermal stability for better cycling performance and safety. Albeit the combined advantages of the conventional electrolytes with LiPF6 and cyclic/linear carbonate solvents, they are not suitable for the high-voltage NCM/NCA materials, due to their easy decomposition and detrimental effects on the cathodes. Novel Li salts such as LiTFSI and LiBOB and organic solvents such as ionic liquids and fluorinated organic compounds are thus developed to replace the traditional Li salts and solvents; however, these Li salts have other issues such as the corrosion of Al current collectors at high voltages and high cost, and the novel solvents show the drawbacks such as high viscosity and poor compatibility with the anodes. The exploitation of novel electrolyte additives with special functions is an economic and effective way to applying the novel Li salts and solvents or to improving the cycling performance and safety of the NCM/NCA-based cells. The SEI formation agents can form thin and stable cathode/electrolyte interface films to protect the cathodes, while the acid scavengers can capture the acid species or H2O in the electrolytes to prohibit the acid corrosion, resulting in the high battery performance. Compared with the organic SEI film formation agents, the Li salt-based inorganic additives can form low-resistance interface films, and the bi-functional additives show the obvious advantage of simultaneous formation of cathode and anode protection films. In addition to the nonflammable solvents, flame-retardant additives are developed to improve the thermal stability and safety of the cells. Especially, the fluorinated phosphates can guarantee the cycling stability and safety of the NCM(A)/graphite full cells, but they are harmful to the environment and human bodies. To further protect the cathodes and electrolytes, overcharge inhibitors are also exploited. Compared to the shutdown-type additives, the redox shuttle additives have outstanding advantages, but many of these redox shuttle additives have negative impacts on the cell performance under the normal operation conditions. The combination of two or more aforesaid electrolyte modification methods will take advantage of the synergistic effects and furtherly make the electrolyte more multifunctional and powerful; however, the compatibilities of the Li salts, solvents and functional additives, and the electrolytes and the electrode materials should be carefully studied. The bi-functional and multifunctional additives are also developed to enhance

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the overall electrolyte/cell performance. The future researches will be focused on the development of bi-/multi-functional additives, and the optimization of electrolyte compositions. The battery safety should not compromise the energy density and cycling stability, and the first-principles calculations will play more important role in designing novel electrolyte materials and understanding their properties. Besides, protic impurities such as water and alcohols have negative effects on the battery performance, and strict control of the preparation and handling of the electrolytes is also indispensable. 6. Summary and outlook Owing to their combined advantages with respect to specific capacity, operation potential, safety, cost and environmental friendliness, NCM/NCA ternary materials are given increasing attention, particularly when the energy density and other property parameters of LiCoO2, LiMn2O4 and LiFePO4 nearly reach the limits. Especially, the Ni-rich materials with low Co contents show greater application prospects in LIBs for EVs, due to their higher energy density and lower cost. Many organizations and companies including U.S. Department of Energy are funding a large number of research groups working on them. To achieve the following goals for their massive LIB applications in long-range EVs (driving range: 300 miles, battery pack: 60−70 kWh, cell level: ~300 Wh kg–1, cathode material level: ~750 Wh kg–1), the NCM/NCA cathode materials should meet the following requirements [16]: (1) specific capacity of ≥200 mAh g–1, (2) cutoff potential of ≥4.4 V (vs. Li/Li+), (3) electrode density of ≥3.9 g cm–3 (currently ≤3.6 g cm–3), and (4) lower cost than LiCoO2 by ≥25%. Albeit the distinct advantages, the NCM/NCA materials are subject to low first-cycle Coulombic efficiency, strong capacity fading, poor rate capability, large voltage degradation, safety issues and low tap/packing densities. By scrutinizing the relationships of chemical composition, crystal structure, particle morphology and properties of the NCM/NCA ternary materials during the different processes, we disclose their performance degradation mechanisms in LIBs, i.e., the changes (or instability) of the materials composition (e.g., Li deficiency/depletion, metal dissolution and oxygen evolution), structure (e.g., cation mixing, phase transition and SEI formation) and morphology (e.g., grain micro-crack and particle rupture), and the corresponding undesirable physical/chemical processes (e.g., side reactions and electric contact loss) mainly during the synthesis and charge/discharge processes hinder the charge transfer (mainly for Li+ ions) and cause the poor performance of the NCM/NCA materials. This review article furtherly summarizes the up-to-date progress made on the NCM/NCA materials with the focus on the materials, electrode and electrolyte designs, and also points out the future research directions. Of course, strict control of the NCM/NCA materials synthesis conditions and the exploitation of novel synthesis methods such as ion exchange and hydrothermal assisted process are also needed to reduce the contents of Li residuals and cation disordering (Fig. 37A, stage 1a).

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Fig. 37 The research directions during (A) NCM/NCA materials preparation and (B) cell fabrication processes.

A series of materials design approaches at the composition, structure and morphology levels are developed to stabilize the cathode materials. The body doping is a simple and effective modification approach to adjusting the chemical composition, and can be operated during the NCM/NCA synthesis process (Fig. 37A, stage 1a). The cation and anion substitution can enlarge the Li slab distance for rapid Li+ ion diffusion, enhance the bond strength between the metal ions and oxygen ions, hinder the Ni migration to the Li slabs, prevent the unwanted phase transition, eliminate the metal and oxygen loss or suppress the side reactions. The doping conditions including the exotic element type, element content and doping method should be carefully adjusted to avoid the detrimental impacts such as the formation of impurity phases and reduction of the total electrode capacity. The simultaneous co-substitution on the Li, transition metal and oxygen sites may offer another avenue to improve the cycling/thermal stabilities. Different from the doping method by the composition change for enhancing the crystal structure, the surface coating with various substances (i.e., core-shell structure

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sometimes) from oxides to fluorides, phosphates, inert Li host compounds, Li+/e– conductors, and cathode materials offers an important structural design approach (Fig. 37A, stage 1b). The coating shells can stabilize the cathode/electrolyte interface, prevent the electrolyte corrosion, hinder the metal dissolution and oxygen evolution, suppress the phase transition or increase the Li+/e– conductivities. Ideal coating layers should be uniformly covered with proper coating thickness, and the coating should not compromise the storage capacity and rate performance. Especially, the structural transition from core-shell to concentration-gradient core-shell and FCG can effectively eradicate the core/shell mismatch, increase the total capacity, prevent the phase change and reduce the side reactions, thus greatly improving the cycling performance and thermal stability; however, further improvement on the preparation method is needed to simplify the synthesis process and cut down the cost (Fig. 37A, stage 3). As another important structural modification method, compositing the cathode particles with a few multifunctional matrices as mechanical and electronic/ionic supports would effectively enhance the particle morphology stability by alleviating the volume change during lithiation/delithiation and facilitate the charge transfer for improving the rate performance. The future research will be focused on the hybrid coating by dual Li+/e– conductive polymers and inorganic composites, concentration-gradient design and the optimization of the preparation process. Beside of controlling the composition and structure of the cathode materials, the morphological modification on the size, pore, shape and configuration of the cathode particles could not only affect the electrochemical/thermal properties but also the tap/packing densities. Especially, the single crystal particles endow potential merits of the structural/morphology/thermal stability, electrode fabrication process and tap/energy densities (Fig. 37A, stage 2), however, the current synthesis methods induce other issues such as the aggregation and wide size distribution of the particles, impurities, complexity in synthesis process, and difficulty in scale-up. Seeking proper synthesis methods for single-particle precursors (without the secondary particle constructure) may offer an important route to obtaining single-crystal NCM/NCA micro-particles. Remarkably, the rational integration of two or more of the abovementioned modification methods should to a great degree enhance the cycling performance on the basis of the synergistic effects. Despite the various modification approaches, the layered NCM/NCA materials must ensure the stability/reversibility of the + – composition/structure/morphology and promote the Li /e transfer during the electrochemical cycling. An additional requirement is to increase the tap/packing densities for high energy/power densities in LIBs. Further improvements on the materials synthesis and modification processes are necessary to achieve the trade-off between energy density, cycle stability, safety and cost.

106

In addition to the materials design at the composition/structure/morphology levels, appropriate electrode design by optimizing the composition and fabrication process also plays an important role in stabilizing the cathodes and facilitating the charge transfer. Small amount of novel conductive nano-additives such as carbon nanotubes and graphene with appropriate size, functional groups and electric conductivity can effectively enhance the cycling performance by fully covering the cathode particles (Fig. 37B, stage 1). Novel binders with high conductivity, high thermal stability and low cost are also needed to improve cycling performance, enhance the cell safety and simplify the battery manufacturing processing. The combination utilization of different conductive additives and binders offers an important route to improve the cell performance. The convenient physical blending of the layered cathode particles with other Li-insert hosts or cathode materials can completely eliminate the initial irreversible capacity loss, improve the thermal stability or lower the whole electrode cost, which are very helpful for their wide applications (Fig. 37B, stage 1). It needs to strictly control the key parameters including the composition, mass ratio, particle morphology, materials interface and treatment condition in order to take full advantage of the properties of each material. Novel or improved electrode fabrication approaches and rigorous processing environment including moisture and dew point should be also investigated for enhancing the overall electrode performance and production efficiency. It is also important to exploit new electrolyte systems to better match the highvoltage NCM/NCA materials for high electrolyte/electrode stability. Despite the special properties of the novel salts such as LiTFSI and LiDFOB, their applications are mainly restricted by the poor dissolubility, corrosion of Al current collectors, or high price. High-voltage solvents such as ionic liquids and fluorinated compounds are also developed, but they have the drawbacks such as high viscosity and poor compatibility with anodes. The utilization of functional additives can facilitate the applications of these novel Li salts and solvents or improve the cycling performance and battery safety. The SEI formation agents and acid scavengers can enhance the NCM/NCA electrode performance by forming stable interface films and prohibiting the acid corrosion, respectively. The Li salt-based SEI formation additives and bi-functional additives exhibit obvious advantages of the resulted low-resistance interface films, and simultaneous generation of cathode and anode protection layers, respectively. In addition to nonflammable solvents, flame retardants and overcharge inhibitors are developed to strengthen the battery safety, but they are detrimental to environment/human bodies and cell cycling performance, respectively. The future researches will be focused on the exploitation of bi-/multi-functional additives and the optimization of electrolyte compositions to furtherly enhance the overall electrolyte properties or solve the tradeoff between cycling performance and cell safety (Fig. 37B, stage 2).

107

Apart from the engineering from the materials, electrode and electrolyte aspects, the whole battery construction must be also well designed to meet the performance and cost targets (Fig. 1). The U.S. Advanced Battery Consortium’s (USABC) targets for 500 km-driving-range EVs are 235 Wh kg–1, 500 Wh L–1 and $125 kWh–1 at battery pack level and 350 Wh kg–1 and 750 Wh L–1 at cell level by 2020 [725]. To reach the driving ranges acceptable to the consumers, the gravimetric and volumetric energy densities of the cells should increase to 220 Wh kg–1 (currently ~180 Wh kg–1) and 300–400 Wh L–1, respectively [128]. A moderate route to cater for these requirements is coupling the high-voltage Ni-rich cathode materials with Si- or Sn-based anode materials (Fig. 37B, stage 3). The utilization of Li metals as anode materials would further reduce the battery cost to <$100 kWh–1, but it is much riskier. The further utilization of highthermostability separators and solid-state electrolytes can greatly enhance the cell safety (Fig. 37B, stage 4), however, other issues such as the low ionic conductivity and high electrode/electrolyte interface resistance needed to be carefully addressed. Another fundamental issue is to figure out the local structure, cycling mechanism and electrochemical changes by combining the first-principles calculations especially based on density-functional theory and molecular dynamics simulations with the advanced characterization techniques such as neutron diffraction, HAADF-STEM, in-situ XRD and EELS. Finally, broader and closer cooperation between the academics and the industrial fields is also indispensable for accelerating the large-scale applications of the NCM/NCA cathode materials in high-performance LIBs. Acknowledgments. This work is partly supported by National High-Tech R&D Program of China (2015AA034601), National Natural Science Foundation of China (91333122, 51402106, 51372082, 51172069, 61204064 and 51202067), Ph.D. Programs Foundation of Ministry of Education of China (20120036120006 and 20130036110012), China Postdoctoral Science Foundation (2018M631419), and Par-Eu Scholars Program, Fundamental Research Funds for the Central Universities (2019QN001). Declarations of interest: none.

108

Table 2 Material, electrode and electrolyte designs for high-performance NCM/NCA cathode applications Modification strategy

Modification approach

Modification type

Sample

Properties or performance

Ref.

Pros and cons

Material composition

Body doping

Cation substitution

Al-doped LiNi0.5Co0.2Mn0.3O2 (in Ni site)

Lower capacity fading of 0.02% per cycle than the undoped (0.07%)

[172]

Ti-doped LiNi0.80Co0.15Al0.02Ti 0.03O2

Initial capacity of 203 mAh g–1 and capacity retention of 74% than the undoped (192 mAh g–1, 67%) after 50 cycles at 1 C 155 mAh g−1 and capacity retention of 84% than the undoped (130 mAh g−1, 69%) after 100 cycles at 1 C

[199]

237 mAh g−1 than the undoped (187 mAh g−1) after 30 cycles at 12.5 mA g−1 between 2.0−4.8 V 122 mAh g−1 and capacity retention of 81% than the undoped (91 mAh g−1, 69%) after 30 cycles at 0.2 C between 3.0−4.3 V 137 mAh g−1 and capacity retention of 83.6% than the pristine (102 mAh g−1, 78.6%) after 100 cycles at 1 C between 2.5–4.8 V 152 mAh g−1 and capacity retention of 90.2% than the undoped (140 mAh g−1, 88.3%) after 50 cycles at 1 C between 2.0–4.8 V High rate capacity (178 mAh g−1 at 10 C) and mid-point voltage retention (95% after 100 cycles)

[185]

Reduce the cathode/electrolyte reactions, enhance the structure stability, inhibit cationic migration and oxygen evolution; undesired LiAlO2 and Al2O3 impurities Render the structure rigid; Li cosubstitutes with Ti on transition metal sites to compensate for charge and results in Li-excess materials Zr4+ can easily enter the Li slabs and function as “pillar” to enhance the crystal structure; excessive doping of Zr could result in impurity phases Occupy Li sites, restrain the Li/Ni mixing and phase transition; high activation energy for Li migration Decrease the Li/Ni disorder, and increase the Li+ diffusion coefficient

Higher initial capacity of ≥160 mAh g–1 and capacity retention of 97% than the pristine (154 mAh g–1, 95%) after 20 cycles at 1 C between 2.6–4.4 V 93.2% capacity retention than the undoped (58.4%) after 100 cycles at 1 C between 2.0–4.6 V 142 mAh g–1 than the undoped (133 mAh g–1) after 50 cycles at 0.5 C between 2.5–4.4 V 283 mAh g–1 and capacity retention of 91% than the undoped (245 mAh g–1, 81%) after 30 cycles at 20 mA g–1 between 2.0–4.8 V 155 mAh g–1 and capacity retention of 97.5% than the undoped (140 mAh g–1, 87.4%) after 200 cycles at 1 C between 3.0–4.3 V

[219]

112 mAh g–1 and capacity retention of

[255]

Zr-doped Li(Ni0.5Co0.2Mn0.3)1– xZrxO2 Mg-doped Li1.2Ni0.12Co0.12Mn0.5 36Mg0.024O2 Ca-doped LiNi0.8– 0.8xCo0.1Mn0.1Ca0.8xO 2

Zn-doped Li1.2Ni0.13– x/3Co0.13–x/3Mn0.54– x/3ZnxO2 V-doped Li1.2Mn0.52– x/3Co0.08–x/3Ni0.2– x/3VxO2 Se6+-doped Li1.2[Mn0.7Ni0.2Co0.1] 0.8–xSexO2

La, Ce and Pr-doped LiNi1/3Co1/3Mn1/3O2

La-doped Li1.2Mn0.54– xNi0.13Co0.13LaxO2 Na+-doped Li1– Na Ni Co Mn x x 1/3 1/3 1/3 O2 K+-doped Li1.20Ni0.13Co0.13Mn0. 54O2 Anion substitution

F–-doped LiNi0.73Co0.12Mn0.15 O2–xFx

(PO4)3−-doped

109

[206]

[192]

[193]

[209]

[217]

Enlarge the interlayer spacing; inactive Zn2+ ions would occupy the Li sites and do not participate in redox reaction Enlarge the interplanar spacing, enhance the structure stability by the strong V-O bond Strongly hybridize with the oxygen ions and reduce the hybridization between Ni and O atoms to enhance the structural stability and impede the O2 evolution Expand the pathway for Li+ intercalation/deintercalation, stabilize the layered framework and suppress the layered-to-spinel transform

[220]

[227]

Enlarge the Li slab space, reduce the cation mixing degree

[230]

Act as fixed pillars in Li layers, large K+ ions can aggravate steric hindrance for the spinel growth

[248]

Increase the lattice parameters, facilitate the primary particle growth, form rock structure on the surface, reduce the O2 release, suppress side reactions Strong covalent bonding combines

LiNi1/3Co1/3Mn1/3O1.9 4(PO4)0.015 SiO44–- and SO42–doped Li-rich cathode materials Cosubstitution

Mg–F co-doped LiNi1/3Co1/3Mn(1/3– x)MgxO2–yFy Al–F co-doped LiNi0.333Co0.333Mn0.29 3Al0.04O2–zFz

Material structure

Surface coating (core-shell structure)

Oxide

Cr2O3-coated NCM

V2O5-coated Li1.2Mn0.54Ni0.13Co0.1 3O2

TiO2-coated LiNi0.6Co0.2Mn0.2O2

ZrOx-coated LiNi1/3Co1/3Mn1/3O2

EPS-coated LiNi0.6Co0.2Mn0.3O2

Fluoride

AlF3-coated Li1.2Mn0.54Ni0.16Co0.0 8O2

Phosphate

AlPO4-coated LiNi0.8Co0.1Mn0.1O2

Li-host compound (ionic conductor)

FePO4-coated LiNi1/3Co1/3Mn1/3O2 LaPO4-coated LiNi0.5Co0.2Mn0.3O2 Al2O3/(AlPO4 or CoPO4) doublelayered Li1.2Mn0.54Ni0.13Co0. 13O2 Li3PO4-coated LiNi0.6Co0.2Mn0.2O2

80% than the undoped (88 mAh g–1, 63%) after 600 cycles at 300 mA g–1 between 2.8–4.5 V SiO44–- and SO42–-doped electrodes showed 220 and 215 mAh g–1 than the undoped (160mAh g–1) after 400 cycles at 30 mA g–1 between 2.0–4.8 V Better cycling stability (180 mAh g–1 after 30 cycles at 20 mA g–1 between 2.8–4.6 V) and thermal stability than the undoped and mono-doped electrodes 150 mAh g–1 and capacity retention of 94.9% than the Al-mono-doped electrode (143 mAh g–1, 84.6%) after 20 cycles at 0.1 C between 3.0–4.3 V 140 mAh g–1 and capacity retention of 83.1% than the uncoated electrode (116 mAh g–1, 72.5%) after 30 cycles at 0.5 C between 3.0– 4.5 V 269 mAh g–1 and capacity retention of 96.3% than the pristine electrode (202.2 mAh g–1, 80.2%) after 50 cycles at 25 Ah g–1 between 2.0–4.8 V

[256]

[262]

[263]

[306]

[285]

Capacity retentions of 85.9% and 80.8% with 186 mAh g–1 than the uncoated electrode (67.5% and 60.8% with 108 mAh g–1) after 100 cycles at 1 C between 2.5–4.3 V at 25 and 55 oC, respectively Capacity retention of 94.6% (2 h ALD coating) than the uncoated electrode (71.3%) after 70 cycles at 1 C between 3.0–4.5 V 93% capacity retention (2 mol% coating) than the uncoated electrode (87%) after 200 cycles at 1 C between 2.8–4.3 V Initial capacity of 208.2 mAh g–1 and capacity retention of 72.4% after 50 cycles at 1 C than the pristine electrode (191.7 mAh g–1, 51.6%) Capacity retention of ~100% than the bare electrode (92%) after 200 cycles at 1 C between 3−4.2 V 143.5 mAh g−1 (2 wt% coating) after 100 cycles at 150 mA g−1 between 2.8−4.5 V 185 mAh g–1 after 60 cycles at 1 C between 3.0–4.8 V 295 mAh g–1 than both the uncoated (253 mAh g–1) and single-layer coated (269–285 mAh g–1) electrodes

[35]

161 mAh g–1 and capacity retention of 94.1% than the uncoated electrode (127 mAh g–1, 76.1%) after 150 cycles at 1 C between 3.0–4.3 V

[393]

110

[273]

[331]

[354]

[138]

[369] [371] [373]

with the MOx (M=transition metal) polyhedrals to stabilize the lattice structure Low Li/Ni cation mixing, excellent layered structure and a few crystal defects Synergistic effects (enhance the crystallinity, increase the lattice parameters, reduce the cation mixing and increase the tap-density) Increase the lattice parameters, suppress the electrolyte decomposition, increase the particle size with increasing F content Inhibit the electrode/electrolyte interface resistance; metal oxides would react with HF to generate metal fluorides and undesired H2O Vanadium ions in 3d0 electronic states can reduce the surface catalytic activity and stabilize the surface oxide ions, VOx coating can suppress the Mn dissolution Protect the cathode from the HF attack, suppress the metal dissolution; react with HF to generate metal fluorides and undesired H2O, poor electronic and ionic conductivities Physically/chemically protect the cathode surface, enhance the Li diffusion, stabilize the electrode/electrolyte interface React with HF, suppress the SEI accumulation, enhance the stability of the cathode surface Eliminate the LiF formation on the particle, inhibit the corrosion by the LiPF6 electrolyte, reduce the metal and oxygen loss Strong P=O bonds can greatly suppress the electrolyte corrosion, strong covalency interactions between the polyanions and metal ions can enhance the thermal stability, metal phosphates can react with the Li residues at high temperatures to generate olivine Li phosphates with high electrochemical and thermal stability

Scavenge both HF and H2O in the electrolyte, high Li+ conductivity for Li+ transport, efficient depletion of Li residuals for formation of high-ion-

Li2Si2O5-coated LiNi0.6Co0.2Mn0.2O2

Initial capacities of 214 and 122 mAh g–1 than the uncoated electrode (197, 92 mAh g–1) at 0.1 and 10 C, respectively, and higher capacity retention of 67% after 200 cycles at 5 C (vs. 52% for the uncoated electrode) Initial capacity of 190 mAh g–1 at 0.1 C and capacity retention of 85% after 50 cycles 164 mAh g–1 and capacity retention of 98% than the pristine electrode (132 mAh g–1 and 87%) after 100 cycles at 1 C between 3.0–4.3 V 165 mAh g−1 and capacity retention 91.1% than the pristine electrode (142 mAh g−1, 78.3%) after 80 cycles at 0.5 C between 3.0−4.5 V 155 mAh g–1 and capacity retention of 84% than the pristine electrode (155 mAh g–1 and 73%) after 50 cycles at 1 C between 2.8–4.6 V 130 mAh g–1 and capacity retention of 94.2% than the pristine electrode (65 mAh g–1, 51.7%) after 50 cycles at 1 C between 2.8–4.5 V 160 mAh g–1 than the pristine electrode (110 mAh g–1) after 100 cycles at 1 C between 3.0–4.5 V

[275]

Ag-coated LiNi1/3Co1/3Mn1/3O2

160 mAh g–1 and retention rate of 94.7% than the pristine electrode (143 mAh g–1 and 86.7%) after 50 cycles at 20 mA g–1 between 2.8–4.4 V

[429]

PEDOT-PEG-coated LiNi0.6Co0.2Mn0.2O2

72 mAh g–1 than the pristine (168 mAh g–1) and PEDOT-coated (160 mAh g–1) electrodes after 100 cycles at 0.5 C; 167 mAh g–1 than the pristine electrode (90 mAh g–1) after 100 cycles at 55 oC 91 mAh g–1 and capacity retention of 90.9% than the Al2O3-coated electrode (83 mAh g–1, 83.1%) after 110 cycles at 10 Ah g–1 between 3.0–4.5 V 205 mAh g–1 than the pristine (190 mAh g–1 after 100 cycles) and 2wt% AlPO4coated (190 mAh g–1) electrodes after 150 cycles at 0.1 C between 3.0–4.8 V 130 mAh g–1 and capacity retention of 95.0% than the pristine (118 mAh g–1, 92.0%) and Li2SiO3-coated (121 mAh g– 1 and 93.7%) electrodes after 90 cycles at 1 C between 3.0–4.3 V 215 mAh g–1 than the pristine electrode (155 mAh g–1) after 50 cycles at 1 C between 2.0–4.8 V

[432]

113 mAh g–1 and capacity retention of

[441]

Li2ZrO3-coated LiNi0.6Co0.2Mn0.2O2 LATP-coated LiNi0.6Mn0.2Co0.2O2

LLTO-coated LiNi0.6Co0.2Mn0.2O2

PTAEP-coated LiNi1/3Co1/3Mn1/3O2

Electronic conductor

Amorphous carboncoated LiNi1/3Co1/3Mn1/3O2 RGO-encapsulated LiNi0.6Co0.2Mn0.2O2

Dual/hybrid conductor

C/Al2O3-coated LiNi0.40Mn0.40Co0.20 O2 RGO/AlPO4-coated Li1.190Mn0.540Co0.143N i0.127O2 Li2SiO3-carboncoated LiNi1/3Co1/3Mn1/3O2

Layered@spinel@ca rbon cathode particle (double shell structure) Cathode

Core-shelled

111

conductivity Li-host compounds

[394]

[403]

[404]

Great enhancement of Li+ ion movement and protection of the electrode from the electrolyte corrosion, excellent storage stability against moisture and air attack

[409]

[426]

[428]

[298]

Facilitate the electron transport, provide electrons to counteract the cation-insertion-induced charge imbalance upon Li+ ion insertion for rapid Li+ ion incorporation, avoid the direct contact with the electrolyte, suppress the oxygen release and detrimental side reactions Increase the electronic conductivity, lower the polarization, promote the formation of protective SEI layer; can dissolve into the electrolyte or be oxidized to silver oxide Provide high Li+ ion and electron conductivities to greatly enhance the charge transfer

Provide high Li+ ion and electron conductivities to greatly enhance the charge transfer; un-uniform coating of the two kinds of conductors

[435]

[439]

[440]

Provide high capacity by the core region, facilitate Li+ ion transport by the spinel interlayer, reduce the HF attack and offer electronic conductivity by the outer carbon layer Combine the high capacity derived

materials

Coreconcentration -gradientshell structure Concentration -gradient core-shell structure

FCG

LiNi0.8Co0.1Mn0.1O2LiNi0.5Mn0.5O2 microparticles Core-shelled LiNi0.80Co0.15Al0.05O2 -LiNi1/3Co1/3Mn1/3O2 microparticles Core-shelled LiNi0.80Co0.15Al0.05O2 -LiFePO4 microparticles Core-shelled LiNi0.5Co0.2Mn0.3O2Li2MnO3 microparticles LiNi0.95Co0.026Mn0.024 O2 (average composition) LiNi0.5Co0.2Mn0.3O2 (average composition)

LiNi0.75Co0.10Mn0.15 O2 (average composition)

LiNi0.85Co0.05Mn0.10 O2 (average composition)

LiNi0.60Co0.15Mn0.25 O2 (average composition) LiNi0.80Co0.06Mn0.14 O2 (average composition)

LiNi0.76Co0.10Mn0.14 O2 (average composition) LiNi0.6Co0.2Mn0.2O2 (average composition) Compositing

Multifunction al electroconductor

LiNi0.8Co0.1Mn0.1O2/ graphene composite

Ionic conductor

LiCoO2/LATP composite granule

98% than the pristine electrode (94 mAh g–1, 81%) after 500 cycles at 1 C between 3.0–4.3 V Capacity retentions of 96.2% and 84.3% than the pristine electrode (86.9%, 68.8%) after 100 cycles at 0.2 C at 25 and 50 oC, respectively Capacity retention of 89% after 400 cycles at 50 oC than the pristine electrode (82% after 250 cycles)

[458]

[459]

215 mAh g–1 after 20 cycles at 0.1 C

[460]

199 mAh g–1 and capacity retention of 90% than the LiNiO2/Li half-cell (161 mAh g–1, 74%) after 100 cycles at 0.5 C between 2.7–4.3 V 189 mAh g–1 than the normal NCM electrode (175 mAh g–1) at 55 oC after 100 cycles at 0.5 C between 3.0–4.5 V

[468]

190 mAh g–1 than the conventional LiNi0.86Co0.10Mn0.04O2 (inner composition) and LiNi0.70Co0.10Mn0.20O2 (outer composition) (118 and 180 mAh g–1) after 100 cycles at 0.2 C between 2.7–4.5 V 190 mAh g–1 and capacity retention of 92% than the conventional LiNi0.80Co0.11Al0.04O2 (150 mAh g–1, 75%) after 100 cycles at 0.5 C between 2.7–4.3 V 185 mAh g–1 and capacity retention of 91% than the conventional half-cell (163 mAh g–1, 83%) after 100 cycles at 0.5 C between 2.7–4.5 V at 55 oC Capacity retentions of 87% and 81.4% than the conventional bulk LiNi0.80Co0.06Mn0.14O2 electrode after 100 cycles at 0.5 C between 2.7–4.3 V at 25 and 55 oC, respectively 197 and 181 mAh g–1 and capacity retentions of 99% and 100% after 100 cycles at C/3 and 1 C, respectively 175 mAh g–1 and capacity retention of 86% than the conventional electrode (133 mAh g–1, 67%) at 1 C after 250 cycles at 60 °C 168 mAh g–1 and capacity retention of 92.2% than the pristine electrode (133 mAh g–1, 76.5%) after 150 cycles at 1 C between 2.7–4.3 V 45 mAh g–1 and capacity retention of 90% after 20 cycles at 0.1 C (in all solid

[471]

112

[463]

[478]

from the core and the high thermal stability stemmed from the shell; the core-shell composition/structure mismatch could result in gradual separation between the core and the shell, inhibit the Li+ ion and electron transfer, and eventually lead to a drastic decline of the battery performance

The concentration gradient within the core or shell would effectively relax the core/shell interface strain that could have formed due to the sharp composition change for improving the long-term cycling performance; the low Mn content in the thin outer shell (≤2 μm ) cannot effectively stabilize the particle surface especially during high-temperature cycling Continuous composition transition from the surface to the bulk enables the FCG cathode to take advantage of high capacity delivered by the Nienriched core and simultaneously the thermal and cycling stability ensured by the Mn-enriched surface layer

[483]

[485]

[486]

[487]

[501]

[395]

Function as mechanical and electrical/ionic supports to alleviate the internal stress resulted from the morphological change of the electrode materials upon cycling and facilitate the transport of charges

state cell) Material morphology

Size

LiMn0.075Co0.775Ni0.15 O2 nanosheets

151.3 and 135.9 mAh g–1 with capacity retentions of 85.1% and 84.7% after 50 cycles at 300 and 3000 mAh g–1 under a high cutoff voltage of 4.4 V, respectively

[512]

Pore/porosity

Porous LiNi1/3Co1/3Mn1/3O2 microsphere

179 mAh g–1 and capacity retention of 90% than the bulk electrode (88 mAh g– 1, 50%) after 50 cycles at 0.1 C between 2.5–4.5 V 200.4 mAh g–1 and capacity retention of 91.2% after 200 cycles at 0.5 C between 2.5–4.5 V 153 mAh g–1 and capacity retention of 82% after 50 cycles

[523]

137 mAh g–1 than the spherical particle electrode (125 mAh g–1) after 100 cycles at 250 mA g–1

[534]

203 and 165 mAh g–1 with capacity retentions of 90% and 88% than the nanoparticle electrode (181 and 122 mAh g–1, 75% and 56%) after 60 cycles at 0.5 and 2 C, respectively 156 mAh g–1 and capacity retention of 82% after 50 cycles at 1 C between 2.5–4.6 V Initial capacities of 192 and 167 mAh g– 1 and capacity retentions of 78.6% and 66.0% after 50 cycles at 0.5 and 10 C between 2.8–4.6 V, respectively 102 mAh g–1 and capacity retention of 78% after 100 cycles at 1 C between 2.8–4.3 V 165 mAh g–1 and capacity retention of 99% after 100 cycles at 0.5 C between 3.0–4.5 V than the Zr-free sample (132 mAh g–1, 78%) 138 mAh g−1 and capacity retention of 78.8% than the NCM (100 mAh g−1, 56.8%) and NCM-Zr (118 mAh g−1, 67.0%) Initial capacity of 203 mAh g−1 at 0.1 C, and capacity retention of 95% after 100 cycles at 0.5 C between 2.7–4.3 V

[529]

The high stability of the hierarchical structure can effectively facilitate the Li transfer and hinder the crack growth during cycling

[90]

Be beneficial to the transportation of Li+ ions along the crystal, reduce the side reactions due to the low surface areas, suppress the micro-crack evolution, enhance the safety in terms of thermal stability; coalescence of the particles into irregular secondary particles and non-uniformity in the shape and size, complex synthesis and subsequent treatment process Stabilize the lattice structure by the Zr doping, reduce the direct contact between the electrode and electrolyte by the Li2ZrO3 coating Synergistic effects of the bulk Zr doping and surface PPy coating

[566]

Synergistic effects of the bulk Al doping and FCG design

Medium-surface-area (68 m2/g) AB led to the highest capacity of 70.5 mAh g−1 and capacity retention of 56.6% of the cells after 100 cycles at 0.1 C within 3.0‒5.1 V

[574]

High electric conductivity, high thermal conductivity, low density, high resistance to acid and alkali, low cost, environmental friendliness; crystallization degree, shape, size, porosity, purity, and surficial defects need to be optimized

Multi-shelled NCM hollow fiber Shape

Hierarchical construction

Single-crystal particle

Doping coating

Co-modification

and

LiNi1/3Mn1/3Co1/3O2 with faceted, edgeblunted polyhedral shape Irregular Li1.2Mn0.54Ni0.13Co0.1 secondary 3O2 particles Nanoporous Li1.2Mn0.54Ni0.13Co0.1 3O2 microsphere

Single-crystal LiNi1/3Co1/3Mn1/3O2 particle Single-crystal LiNi0.6Co0.2Mn0.2O2 particle Single-crystal LiNi1/3Co1/3Mn1/3O2 particle Zr-doped LiNi1/3Co1/3Mn1/3O2 coated with Li2ZrO3 Zr-doped LiNi0.5Co0.2Mn0.3O2 coated with PPy

Doping FCG

Electrode composition

Conductive additive

and

Conventional additive

Al-doped FCG LiNi0.76Co0.09Mn0.15 O2 (average composition) Acetylene black (AB) and carbon black (CB)

113

[524]

[533]

[548]

[36]

[374]

[563]

Small particles with high specific surface area and short diffusion distance can facilitate Li transfer; too small particles induce more surface reactions, the high reactivity of the nanoparticles is detrimental to safety, reduce the tap and energy densities Increase the contact area with electrolyte, shorten Li+ diffusion path, alleviate the volume change upon cycling, inhibit the particle cracking; porous particles are easy to crack during the electrode pressing period Enhance the mobility of charge carriers, alleviate the lattice strain during cycling, affect the tap density

Novel additive

Carbon fiber

Li1.18Ni0.15Co0.15Mn0.52O2/Li cells with 3 wt% carbon fibers and 12 wt% CB showed the highest capacity of 239 mAh g−1 and capacity retention of 94.4% after 50 cycles at 40 mAh g−1 within 2.0‒4.8 V

[577]

MWCNT

Enhanced rate capability at 0.25–5 C (e.g., 87 and 58 mAh g–1 at 5 C for MWCNT- and carbon black-based LiNi1/3Co1/3Mn1/3O2 electrodes, respectively) Graphene/carbon black-containing LiNi1/3Co1/3Mn1/3O2/Li cells showed the best rate capacity of 78 mAh g−1 at 2000 mAh g−1 within 2.5‒4.1 V

[582]

P3HT-CNT modified LiNi0.8Co0.15Al0.05O2 electrode showed improved capacity of 80 mAh g−1 (80% of the initial capacity) after 1000 cycles at 16 C, a value that was four times greater than that obtained in the PVDFbased electrode

[590]

Fluorinated polyimide (FPI)-based Li1.2Ni0.132Co0.13Mn0.54O2/Li cell displayed higher capacity of 223‒198 mAh g−1 and capacity retention of 89% than PVDF-based cells after 100 cycles at 0.2 C and 55 oC within 2.5‒4.7 V CMC-based LiNi1/3Co1/3Mn1/3O2/Li cathodes had higher capacity of 141.9 mAh g−1 and capacity retention of 90.1% than alginate- and PVDF-based electrodes (126 mAh g−1 and 89.2%, and 111.7 mAh g−1 and 86.3%) after 200 cycles at 0.5 C within 2.5‒4.6 V and also better rate performance High initial discharge capacity of ∼280 mAh g−1 with no irreversible capacity loss and excellent cyclability

[592]

High thermal stability and the enhance cycling performance and cell safety at high temperature; maybe detrimental to the room-temperature cell cycling performance

[601]

Better cycling performance than the Li1.1Mn1.9O4 and LiNi0.4Co0.3Mn0.3O2 cathodes after 30 day storage at 45 oC Average discharge capacity of the

[617]

Low cost, environmental friendliness, simplification of the battery manufacturing; strong reactivity between the metal oxide materials and water results in the leaching of Li+ and metal ions from the cathodes and the increase of pH that in turn causes the corrosion of Al current collector Li-insertion hosts can recapture the Li+ ions that cannot be re-intercalated into the cathode structures during the discharge process, instead of solving the underlying causes of the metal ion migrations and oxygen loss; reduce the active cathode materials loading Mixing two cathode materials could circumvent the shortcomings of the parent materials to achieve higher energy or power density as well as

Graphene

Binder

Conventional binder

PVDF

Conducting binder

Poly(3,4ethylenedioxythioph ene) (PEDOT), polypyrrole (PPy), poly(3hexylthiophene-2,5diyl) (P3HT) and polythiophene (PT) Polyamide-imide (PAI) and polyimide (PI)

High-thermostability binder

Blending

Water-soluble binder

Carboxymethyl cellulose (CMC), alginate, guar gum, and polyacrylic latex (PAL)

Li-host material

Li1.2Mn0.54Ni0.13Co0.1 blended with 3O2 Li4Mn5O12 or LiV3O8

Cathode material

LiNi0.4Co0.3Mn0.3O2 blended with spinel Li1.1Mn1.9O4 LiNi0.80Co0.15Al0.05O2

114

[585]

[609]

[618]

Low density, high bulk conductivity, high thermal conductivity, and enhanced mechanical stability of the electrodes by the one-dimensional fibers; complicated synthesis process and high cost of the carbon fibers, and the uncompleted coverage of the cathode particles by the fibers Partial use of MWCNTs (0.5–2 wt%) with high conductivity and crosslinked charge-transfer paths in the electrode can improve the energy and power densities and cycling stability Appropriate size, surface chemistry and electrical conductivity of the graphene additives can greatly improve the cell performance Wide electrochemical window, high adhesivity with electrode materials, high viscosity of the organic slurry, low swelling ratio and dissolubility in electrolyte; nonconductive to Li+/e‒, much expensive, and use of organic solvents such as toxic and lowvolatility NMP High electro-conductivity, electrochemical stability, and easy integration into electrode structures; high cost

blended with spinel Li1.1Mn1.9O4

NCM and LiFePO4 mixture

Electrode fabrication

Vacuumfiltering

Freestanding, binder-free cathode

LiNi0.5Co0.2Mn0.3O2/ CNT flexible film cathode

Electrolyte composition

Li salt

Conventional Li salt

LiPF6

Novel Li salt

LiTFSI, LiFAP, LiDFOB EC

Solvent

Conventional solvent

LiFSI, LiBETI,

DEC, DMC, EMC Novel solvent

Li1.1Mn1.9O4 and LiNi0.80Co0.15Al0.05O2 electrodes but greatly improved capacity retention even at 60 oC, expansion in the operating voltage range Continuous variation of voltage appeared in the cell voltage profile; improvement in capacity retention, rate performance and thermal stability Reduce the electrode polarization, and lead to a ~25% increase in the reversible capacity and an obvious improvement in cycling stability and rate performance, compared to the conventional slurrycoated electrode See Table 3 for details

enhanced thermal stability and lower cost

[604, 621]

[629]

See Table 3 for details

See Table 4 for details

See Table 4 for details

Ionic liquids

Fluorinated organic solvents Novel electrolyte system

Highconcentration system

8.67 mol/kg LiBF4/DMC

EC-free solvents with additives

EMC/FEC

LiPF6:FEC/HFDEC/ 1% LiDFOB

Additive

SEI formation agent

N-allyl-N,N-bis(trimethylsilyl)amin e (NNB) Fluoroethylene carbonate (FEC), ES and VC

High-concentration electrolyte-based LiNi0.8Co0.1Mn0.1O2/Li cells exhibited higher capacity retention of 93.3% than the conventional electrolyte-based cells (i.e., 1 M LiPF6:EC/DMC) after 50 cycles at 0.1 C within 3.0‒4.3 V LiNi0.4Co0.2Mn0.4O2/graphite cells with 98% EMC and 2% FEC showed an increased energy density by at least 10% with a cutoff voltage of 4.4 V Novel electrolyte-based LiNi0.5Co0.2Mn0.3O2/graphite cells displayed a higher capacity retention of 82% than the conventional electrolytebased cells (67%) after 100 cycles at C/3 within 3.0‒4.6 V Capacity retention of the NCA electrodes increased from 72.8% to 86.2% after 300 cycles at 1 C within 3.0‒4.2 V 2 vol% FEC-containing LiNi1/3Co1/3Mn1/3O2 electrode showed the highest capacity of 114 mAh g−1

115

[642]

[647]

Enlarge the voltage variation range for effectively monitoring the SOC of the LiFePO4-based battery, improve the electrochemical/thermal properties Further improve the battery performance; un-scalable electrode preparation process and high cost

High conductivity, excellent roomtemperature performance, and low cost; decomposition with water and poor properties at high/low temperature Wide electrochemical window, high thermal stability; corrosion of Al foils at high voltage, or high cost High permittivity, formation of stable anode interface films; high viscosity, narrow electrochemical window Low viscosity; low permittivity, and poor thermal stability Wide electrochemical window, low volatility and high flame retardancy; high viscosity, low conductivity, low wettability and the resulted poor cell performance High thermal stability, elevated charge potential; high viscosity and incompatibility of the graphite anode Wide electrochemical window, inhibition of the electrolyte decomposition; high viscosity

Improve the cell performance by elevating the cutoff voltage and forming stable SEI films

[652]

[680]

[682]

Decomposition or polymerization before the baseline electrolytes and form stable SEI films to protect the cathodes; the resulted lowconductivity SEI films may increase the cathode/electrolyte interface resistance

after 150 cycles at 1 C Dopamine

Lithium difluorophosphate (LiDFP)

LiDFBOP

Acid scavenger

Diphenyldimethoxys ilane (DPDMS)

1-(2-cyanoethyl) pyrrole (CEP)

Flame retardant

Over-charge inhibitor

Dopamine-contained LiNi1/3Co1/3Mn1/3O2/graphite cells exhibited higher capacity of 147 mAh g−1 and capacity retention of 90.1% than the dopamine-free cells (136 mAh g−1 and 83.3%) after 100 cycles at 1 C within 3.0‒4.5 V 1 wt% LiDFP-containing cells displayed higher capacity of 177.0 mAh g−1 and capacity retention of 90.2% than the 1 wt% PES-containing (161.0 mAh g−1 and 77.1%) and additive-free (113.1 mAh g−1 and 64.0%) cells after 200 cycles at 1 C within 3.0‒4.45 V LiDFBOP-containing LiNi0.5Co0.2Mn0.3O2/graphite pouch cells exhibited higher capacity retentions (75% and 93%) and Coulombic efficiencies than the additive-free cells (21% and 73%) at both 25 and 0 oC 1 wt% DPDMS-containing LiNi0.6Co0.2Mn0.2O2/Li cells showed a higher capacity retention of 93.3% than the DPDMS-free cells (71.9%) after 200 cycles at 2 C and 55 oC within 2.8‒4.3 V 1 wt% CEP-containing LiNi0.6Co0.2Mn0.2O2/graphite cells showed higher capacity retention (81.5%) than the CEP-free cells (27.4%) after 50 cycles at 1 C within 3.0‒4.35 V

Alkyl phosphorusbased compounds such as trimethylphosphate (TMP) and triethylphosphate (TEP) Fluorinated phosphates such as tris-(2,2,2trifluoethyl) phosphate (TFP), and bis(2,2,2trifluoroethyl) methyl phosphate (BMP) Shutdown-type additives

Redox shuttle additives such as 1,3dimethylimidazolidi

[684]

[688]

Formation of high-conductivity interface films on cathodes from the Li salt decomposition to protect the cathodes and facilitate Li ion transfer

[695]

Simultaneous formation of highconductivity cathode and anode interface films to inhibit the electrolyte decomposition, protect the cathode and anode from destruction, and promote the Li+ ion transport Scavenge HF and PF5 in the electrolytes, protect the cathodes from electrolyte corrosion

[704]

[705]

Prevent HF generation by eliminating H2O, and inhibit the transition metal dissolution from HF corrosion

[643, 709, 712, 713]

Generate free radicals during the electrolyte combustion process and suppress the liberation of flammable gas; instability at low potentials and ineffectiveness to form stable SEI films, high viscosity and environmental pollution Guarantee the high electrolyte conductivity, form stable SEI films on anodes, and exhibit high flame resistance; the halogen-based flame retardants are harmful to environment and human bodies

[714716]

[643]

NHC-BF3 and NHC-PF4CF3 were identified as promising overcharge additives for LiNi1/3Co1/3Mn1/3O2 electrodes up to 4.4 and 4.6 V vs. Li/Li+,

116

[722]

Polymerization at certain potentials and form insulating polymer films on the cathodes to prohibit overcharge; irreversibility of the shutdown process and the potential increase of the internal pressure in the cell related to the polymer layer formation Reversible oxidization at defined potentials on the cathode side and reduction at anodes to prohibit cell overcharge; detrimental to cell

Multifunction al additive

n-2-mmtrifluoroborate (NHC-BF3) and NHC-PF4CF3 Diethyl(thiophen-2ylmethyl)phosphona te (DTYP)

respectively

0.5 wt% DTYP-based cells showed higher capacity retention of 84% than the additive-free cells (18% ) after 280 cycles at 1 C and 60 oC within 3.0‒4.9 V, due to the formation of stable SEI films and effective inhibition of HF corrosion; the DTYP-based electrolyte showed higher thermal decomposition onset temperature of 223 oC than the baseline electrolyte (193 oC)

performance due to the formation of high-resistance SEI films from the additive decomposition [724]

Improve the overall properties of the electrolytes/cells with the synergistic effects of two or more functional additives; tradeoff between the cell performance and safety

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CVs:

Lehao Liu is a postdoctoral researcher in North China Electric Power University. He received his PhD degree in 2016 from Northwestern Polytechnical University under the supervision of Prof. Tiehu Li. During 2012–2014, he worked as a joint-training PhD student at University of Michigan with the guidance of Prof. Nicholas A. Kotov. After his PhD graduation, he began to work as a chief engineer in Citic Guoan MGL Power Source Technology Company in Beijing. His research interest is focused on the preparation of nano-/micro-materials for Li-ion battery applications. He has published more than 20 peer-reviewed papers in the field.

Meicheng Li is a professor in North China Electric Power University. He is also the Director of New Energy Materials and PV Technology Center, and the Vice Dean of the School of Renewable Energy. He obtained his PhD degree at Harbin Institute of Technology in 2001. He worked in University of Cambridge as Research Fellow from 2004 to 2006. He won the Excellent Talents in the New Century by the ministry of education in 2006. His current research topic is the New Energy Materials and Devices, such as solar cells and lithium ion battery. He has contributed more than 200 journal articles. He got almost more than 10 items of awards for the science and technology success. He is an executive fellow of the China Energy Society, fellow of Chinese Society for Optical Engineering.

1

Lihua Chu is a lecturer in North China Electric Power University. She received her PhD degree in Condensed Matter Physics in 2013 at Beihang University and Bachelor degree in Material Science and Technology in 2008 at Central South University. She was a postdoctoral fellow at North China Electric Power University during 2013-2015. During 2011-2012, she worked as a combined training of postgraduate at Centre national de la recherche scientifique (CNRS), France. Her current research focuses on the preparation of nanomaterials and its application in energy field, and structural functional materials.

Bing Jiang is an associate professor of in State Key Laboratory of Alternate Electrical Power System with Renewable Energy Sources, North China Electric Power University. She received his PhD degree of Material Physics and Chemistry from University of Science and Technology Beijing in 2009. Currently, her primary research interest is fabrication and characterization of nanomaterials and nanostructures that are applied in renewable energy-related devices, such as solar cell and lithium ion battery.

Ruoxu Lin is a senior engineer in Citic Guoan MGL Power Source Technology Company in Beijing. He received his PhD degree in Material Physics and Chemistry from Beihang University in 2016. His main research interest is focused on the synthesis of lithium nickel cobalt manganese/aluminum oxides cathode materials and their battery application in electric vehicles and portable devices. He has published more than 15 papers in the Li-ion battery field.

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Xiaopei Zhu is a senior engineer in Citic Guoan MGL Power Source Technology Company in Beijing. She received her Master degree in University of Science and Technology Beijing in 2009 and then joined the Citic Guoan MGL Electric Power Sources Company. She is now pursuing a PhD degree in University of Science and Technology Beijing with the guidance of Prof. Lizhen Fan. Her research work is focused on Ni/Co/Mn-based cathode materials for Li-ion battery applications. She has obtained over 20 patents about the cathode materials.

Guozhong Cao is Boeing-Steiner Professor of Materials Science and Engineering and Adjunct Professor of Chemical and Mechanical Engineering at the University of Washington. He received his PhD degree from Eindhoven University of Technology in Netherlands. He has published over 300 refereed papers, authored and edited 7 books, and presented over 200 invited talks and seminars. He serves as editor of Annual Review of Nano Research and associate editor of Journal of Nanophotonics. His current research is focused mainly on nanomaterials for energy related applications, such as lithium-ion batteries, solar cells, supercapacitors, and hydrogen storage.

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