Identifying and Addressing Critical Challenges of High-Voltage

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Identifying and addressing critical challenges of highvoltage layered ternary oxide cathode materials Shu Zhang, Jun Ma, Zhenglin Hu, Guanglei Cui, and Liquan Chen Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.9b01557 • Publication Date (Web): 25 Jul 2019 Downloaded from pubs.acs.org on July 27, 2019

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Chemistry of Materials

Identifying and addressing critical challenges of high-voltage layered ternary oxide cathode materials

Shu Zhang,† Jun Ma,*,† Zhenglin Hu,†,§ Guanglei Cui,*,† and Liquan Chen†,‡ † Qingdao Industrial Energy Storage Research Institute, Qingdao Institute of Bioenergy and Bioprocess Technology, Chinese Academy of Sciences, Qingdao 266101, P. R. China § Center of Materials Science and Optoelectronics Engineering, University of Chinese Academy of Sciences, Beijing 100049, P. R. China ‡ Key Laboratory for Renewable Energy, Beijing Key Laboratory for New Energy Materials and Devices, Beijing National Laboratory for Condensed Matter Physics, Institute of Physics, Chinese Academy of Sciences, Beijing 100190, P. R. China ABSTRACT Layered ternary oxide cathode materials LiNixMnyCo1-x-yO2 (NMC) and LiNixCoyAl1-x-yO2 (NCA) (referred to as ternary cathode materials, TCMs) with large reversible capacity, high operating voltage as well as low cost are considered as the most potential candidate materials for high energy density lithium ion batteries (LIBs) used in hybrid electric vehicles (HEVs) and electric vehicles (EVs). However, next-generation long-range EVs require an energy density of 800 W h kg-1 at the cathode level, which cannot be obtained using the commercially available TCMs. Developing highvoltage TCMs is a promising solution to enhance energy density of LIBs. Nonetheless, the capacity decay, poor long-term cycle life and microcrack at high operating voltage have limited their practical applications. In this paper, the development of the high-voltage TCMs is reviewed from degradation mechanism, cathode electrode modification, electrolyte design, solid state electrolytes and so on. The critical factors, recent progress and perspectives that improve the performance of TCMs with high-voltage operation are reviewed, which could provide important information and precautions to the practical use of these cathode materials under high operating voltage. 1. INTRODUCTION Lithium ion batteries (LIBs) as the rechargeable power sources with high capacity and high operation voltage outperform other available battery systems in terms of energy density, design flexibility, cycle life, and low self-discharge rate 1. Since the introduction of rechargeable LIBs to the consumer market in the early 1990s, they have accelerated global communication with the revolution in portable electronics. Nowadays, LIBs have also shown extensive application prospects in electric vehicles (EVs), hybrid EVs (HEVs) and smart grids 2, 3. However, compared with internal

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combustion vehicles, recent EVs and HEVs suffer from range anxiety due to the limited energy density of power LIBs 4-9. Therefore, it is necessary to sustain the research and development efforts to improve the energy density of power LIBs. LIBs have four key components: cathode, anode, electrolyte and separator. The output voltage of the battery is determined by the difference in electrochemical potentials between the cathode and anode 10. In general, the output voltage is determined by the cathode because graphite is the most used commercial anode. Furthermore, cathode is the provider of lithium ion, which is also called “lithium sources”, so the capacity of LIB mainly depends on the cathode materials. Thus, the design of appropriate cathode systems is of prime importance to realize a high energy density (capacity  voltage) LIB. So far, the cathode materials can be classified into three categories: 1. Lithium oxide compound with layer structure LiMO2 (M = Co, Ni, Mn), such as LiCoO2 11, LiNixCoyMn1-x-yO2 12 2. Spinel structure compounds. The typical materials are LiMn2O4

14,

3. Polyanionic structure compounds. The typical materials are LiFePO4

16,

and LiNixCoyAl1-x-yO2 LiMn1.5Ni0.5O4

15;

13;

LiFexMn1-xPO4 17 , Li3V2(PO4)3 18. By comprehensive consideration of the factors of energy density, cost, security, and preparation technical maturity, LiFePO4 and layered LiNixMnyCo1-x-yO2 (NMC) and LiNixCoyAl1-x-yO2 (NCA) (referred to as ternary cathode materials, TCMs) are the most suitable cathode materials for power LIBs. LiFePO4 displays high safety characteristic and cycling performance, but the low working voltage and low volume specific energy limit its applications in long-range EVs 11, 14, 19-23. In sharp contrast, layered TCMs have much larger specific capacity, rate capability, and working voltage than bare LiFePO4, and have already been widely applied in longrange EVs 24-28. For instance, Tesla EVs are the most successful model of NCA cathode materials applied in power LIBs. However, next-generation long-range EVs require an energy density of 800 W h kg-1 at the cathode level, which cannot be obtained using the current commercial TCMs 29. The main strategy for extracting a higher energy density from TCMs to meet the drive range requirement for EVs is to increase the capacity and operating voltage following the equation: energy density = capacity  voltage. In general, it is believed that the predominant oxidation states of Ni, Co and Mn in the NMC compound are +2, +3 and +4 respectively with small content of Ni3+ and Mn3+ ions 30. For NCA materials the oxidation states of Ni, Co and Al are all +3. The crystal field produced by octahedral oxygen anion coordination splits the five d orbitals of the transition metal(TM) into two sets of


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levels labeled eg and t2g. The eg levels have a higher energy than the t2g levels due to the increased electrostatic repulsion from the coordinating anions. The commonly observed electronic configurations of d orbitals in cationic Co, Mn, and Ni species are shown in Figure 1a and b. Figure 1c qualitatively summarizes the positions of the redox active energy levels for Ni, Mn, and Co, as well as the oxygen 2p band in layered LiMO2 (M = Co, Ni, Mn). The Co3+/Co4+ level is the lowest in energy (and therefore highest in voltage) because Co redox involves the more stable t2g orbitals, while Ni and Mn redox involves the less stable eg orbitals. The next highest energy levels above those of Co are the Ni3+/Ni4+ and Ni2+/Ni3+ couples followed by that of the Mn3+/Mn4+ redox couple. Thus, the deintercalation of a typical layered material proceeds via the oxidation of Ni2+ to Ni3+, followed by the oxidation of Ni3+ to Ni4+, and then finally the oxidation of Co3+ to Co4+. The Mn4+ remains electrochemically inactive throughout the entire process. Therefore, the capacity improvement of TCMs mainly depends on two factors: Ni content and operating voltage. The increase of Ni content in TCMs under low operating voltage ( 4.3 V) for the oxidation of Ni2+ to Ni3+/Ni4+ contributes the main capacity. Higher operating voltage ( 4.3 V) also can enhance the capacity due to the oxidation of Co3+.

Figure 1. (a) Crystal field spitting of d orbitals in an octahedral environment. (b) Typical electron configurations in layered oxides. (c) Qualitative positions of energy levels in layered oxides. Reprinted with permission from ref 31.

Copyright 2017 Wiley.

However, TCMs with high Ni content or high operating voltage suffer a rapid capacity decay and limited service life 32-35. The main opinion on the capacity fading of Ni-rich TCMs is that high Ni content results in severe cation mixing. Consequently, the cation mixing leads to phase transformation from well-ordered layered phase to spinel and/or rock salts phase. Besides, the highly oxidizing Ni4+ ions promptly react with the organic electrolyte on the electrode/electrolyte interface, which result in the surface structure transformation, electrolyte decomposition and O2 release. Recent years, researchers have proposed many strategies to solve these problems in Ni-rich TCMs,



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such as surface modification, element doping, electrolyte additives, and so on. Some reviews have discussed the progress of Ni-rich TCMs in detail

29, 36-39.

However, there is a lack of in-depth

discussion focusing on TCMs issues with high-voltage operation. Significantly, high operating voltage will exacerbate phase transformation and electrode/electrolyte interface reactions of TCMs during cycling, especially in Ni-rich TCMs. Furthermore, high operating voltage brings about large volume change and even structural microcracks, which will lead to contact loss and recurring interface reactions. Therefore, in this review the failure mechanism, modification, remaining challenge and the perspectives of high-voltage TCMs are detailed discussed, which will provide important information and precautions to the development of TCMs under high operating voltage and have great significance to the research in basic science and industrial production. At the end of this review we put forward some suggestions for future research of the high-voltage TCMs-based LIBs, which can promote the development of their practical use in EVs. 2. DEGRADATION MECHANISMS Degradation is mainly characterized by irreversible capacity loss and voltage drop, coupled to changes in kinetic and thermodynamic properties of involved materials and their interfaces 40. Many degradation phenomena have to be considered on loss of contact between particles, cracking of particles, point defect formation as well as structural changes inside the active materials and interface reactions (Figure 2). In the following, we will discuss the origin of aforementioned degradation phenomena from the aspects of cation mixing, phase transformation, interface reaction, and microcrack of high-voltage TCMs. The dominant factors to capacity degradation of Ni-poor and Ni-rich TCMs may be different to some extent between the bulk and surface, but not different in essence and both of them refer to the influence factors we mentioned above. So we discuss the degradation mechanisms at high voltage for all the TCMs without distinguishing the Ni-poor and Ni-rich.


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Figure 2. Overview on basic degradation mechanisms of high-voltage TCMs. CEI is the abbreviation of cathode electrolyte interface which is the reaction product of cathode and electrolyte.

2.1 Cation mixing and phase transformation

Figure 3. Illustration of the ordered and disordered phase in layered lithium metal oxides and their structural _

_

transformation. (a) Well ordered R 3 m structure; (b) The cation disorder or cation mixing phase with Fm 3 m _

structure; (c) R 3 m structure with Li vacancies in highly charged state; (d) Partially cation mixed phase with TM



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ions in Li slab. Reprinted with permission from ref 36. Copyright 2015 Wiley.

Figure 3 shows the structure of the layered materials and their structural transformation36. Generally, LiMO2 (M = Co, Ni, Mn, Al, etc.) cathode materials have a layered NaFeO2 structure _

with R 3 m space group with alternating layers formed by edge-sharing LiO6 and MO6 octahedra (Figure 3a). The Li ions are on 3a sites, the M ions are randomly placed on 3b sites, and oxygen atoms are on 6c sites

41 31.

The M ions and Li ions that occupy the octahedral sites of alternating

layers with an ABCABC… stacking sequence along the rhombohedral [001] direction known as“O3-type” structure

38.

And during charging process the O3 phase may transform to O1 phase

with AB stacking. By tracking the variations of lattice parameters, the conversion of hexagonal phase H1 to hexagonal phase H2 and hexagonal phase H3 can be observed upon the deep-degree delithiation and lithiation processes 42, 43. The H1, H2 and H3 phase show the same hexagonal phase. H1 and H2 are O3 phase and H3 is O1 phase 28. They have some difference in lattice parameters which is attributed to the extraction/insertion of lithium ions. Similar variations can be observed in layered LiNiO2 systems or other layered O3 type compounds cycled in high potential regions 44, 45. These phase transformations are detected in the bulk structure and always confirmed by XRD. Because it is difficult to distinguish between layered phase H3 and cubic spinel by synchrotron radiation diffraction, the cubic spinel nano-domains may present. The formation of the layered phase H3 or/and spinel phase at higher charged voltage causes the collapse of the layered host structure in the c direction, resulting in an irreversible structural transformation and thus capacity fading 46. Owing to the similarity of ion size, Li+/Ni2+ cation mixing involved cations disordering between _

M ions sites and Li ions sites and LiMO2 transformed to Fm 3 m structure (Figure 3b). The cation mixing would attract oxygen ions from adjacent oxygen layers, locally shortening the oxygenoxygen distance between neighboring MO2 slabs. As a result, the smaller distance between the slabs leads to a higher activation energy barrier for lithium diffusion. And the lithium diffusivity is lower because of the hindrance caused by the transition metal in the lithium layer, resulting in relatively poor electrochemical performance. This cation migration from 3b to 3a site occurs not only during the material synthesis procedure but also the electrochemical cycling process, especially more serious at a high cut-off voltage 47, 48. In a highly charged state, the material has an unstable structure owing to the increase in the extraction of Li from the structure, and this instability leads to transition-


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metal-ion migration from the transition-metal layer to the lithium layer. The transition metal ion migration has a tendency to occupy every second Li vacancy because the disordered cations repel each other (Figure 3c, d). The probable reason of the cation mixing is widely accepted as the similar ionic radius of Li+ and Ni2+. Nevertheless, Doeff et al. reported that the cation mixing layer consisted of not only Ni2+ but also of other metal ions, such as Mn2+ and Co2+ 49. Bi et al. found that the cation mixing also related to the oxygen deficiency 50. The oxygen defects should trigger Ni/Li exchanging positions in the bulk, and accelerate the decomposition on the interface and surface. The increase of oxygen deficiency significantly strengthened the aforementioned instabilities, which severely deteriorate the electrochemical performance due to the seriously aggravated polarization. The layered oxides materials exhibit strong thermodynamic driving forces upon deintercalation to transform from well-ordered layered phase to spinel via the cation mixing and to transform to rocksalt via the evolution of oxygen during cycling 48. The spinel and disordered rocksalt structures are crystallographically related to the O3 structure in that they all have an fcc oxygen framework, differing only in their arrangement of M cations over the octahedral and tetrahedral interstitial sites. One quarter of the metal ions migrate from the transition-metal layer to particular octahedral sites in the Li layer, resulting in metal ions ordering within the fcc oxygen framework and cubic symmetry, leading to the transformation from O3 to spinel. The rocksalt structure represents a disordering of the cations in O3 structure with the release of oxygen, while the Li and transitionmetal layers no longer exhibit long-range periodic order. Compared to the well-ordered phase, the spinel and/or rock-salt structures possess higher activation energy barrier for Li+ diffusion owing to the reduced distance between the slabs. Therefore, the disordered spinel and/or rocksalt structures with TMs in the lithium layer exhibit a lower Li+ diffusion coefficient and sluggish Li+ kinetics during redox reactions, finally leading to the capacity fade and poor rate performance of TCMs. An exceptionally oxidative environment imposed by the high-voltage charge is likely to trigger a severe oxygen evolution from the electrode and induce phase transformation. Transformations to the cubic phases nevertheless occur near the surface, having a rocksalt-like phase with approximate composition MO on the surface and a spinel-like phase in the subsurface undergoing high-voltage cycling conditions

51.

Upon cycling and/or electrolyte exposure, the surfaces of NMC particles



_

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_

undergo reconstruction from an R 3 m layered structure to an Fm 3 m rock-salt structure in addition to the buildup of a complicated surface reaction layer containing LiF and organic components. Interestingly, the structural reconstruction at the surface is highly anisotropic, and primarily occurs along the lithium ion transport direction (Figure 4). While the layered structure of the O3 phase can be seen in the bulk, the presence of transition-metal ions in the Li layer (characteristic of the cubic phases) can be seen near the surface. This suggests that protection coatings on certain orientations of NMC particles should be more beneficial than those on other orientations.

Figure 4. Atomic resolution ADF-STEM images of NMC particles. (a) After electrolyte exposure (the exposure time is ~30 hours, equivalent to the time used for one full cycle in this study). (b) After 1 cycle (2.04.7 V); the blue arrow indicates the surface reconstruction layer. (c, d) FFT results showing the surface reconstruction layer (Fm-3m [110] zone axis) and the NMC layered structure (R-3m [100] zone axis), respectively in (b). (e) Showing variation of the surface reconstruction layer thickness on orientation after 1 cycle (2.04.7 V). (f) Image showing loose atomic layers on an NMC particle, after 1 cycle (2.04.7 V). The blue lines indicate the boundaries between the NMC layered structure and surface reconstruction layer in all images. The scale bars are 2 nm in all images. Reprinted with permission from ref 49. Copyright 2014 Macmillan Publishers.

The relation between surface structure evolution and oxygen vacancy of high-voltage TCMs has also been explored. Sooyeon Hwang et al. reported the local evolution of the surface structure of


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LixNi0.8Co0.15Al0.05O2 as a function of state of charge (SOC) (0.1 ≤ x ≤ 0.5) using transmission electron microscopy (TEM) and electron energy loss spectroscopy (EELS) (Figure 5)

52.

It was

_

found that the surface changes from the layered structure (space group R 3 m) to the disordered spinel _

_

structure (Fd 3 m), and eventually to the rock-salt structure (Fm 3 m), and that these changes are more substantial as the extent of charge increases. EELS results indicate that these crystal structure changes are also accompanied by significant changes in the electronic structure, which are consistent with delithiation leading to both a reduction of the Ni and an increase in the effective electron density of oxygen. This leads to a charge imbalance, which results in the formation of oxygen vacancies and the development of surface porosity. The degree of local surface structure change differs among particles, likely due to kinetic factors that are manifested with changes in particle size.

Figure 5. (a) High-resolution image of an overcharged LixNi0.8Co0.15Al0.05O2 surface region. The overall particle is shown as an inset. (b) Selected area diffraction patterns from whole particle, and (i−iii) fast Fourier transformation results of three regions in (a), respectively. L, S, and R denote the layered, spinel, and rocksalt structures, respectively. (c) Schematic summarizing the crystallographic and electronic structure changes that occur in LixNi0.8Co0.15Al0.05O2 cathode material upon electrochemical charging. Reprinted with permission from ref 52. Copyright 2014 American Chemical Society.

2.2 Interface reactions Besides the intrinsic structural stability of TCMs, the interface stability between TCMs and other battery components is also essential for obtaining excellent performances for LIBs. The open circuit voltage (OCV) is the difference in chemical potential between the anode and cathode. This working



voltage is also limited by the electrochemical window of the electrolyte, which is determined by the energy gap from the highest occupied molecular orbital (HOMO) to the lowest unoccupied molecular orbital (LUMO)

53.

But the HOMO and LUMO describe the electronic properties of

isolated molecules, and their energy levels do not indicate species participating in redox reactions. To more clearly describe the electrochemical stability of electrolytes, Peljo et al. have provide a thermodynamic representation based on redox potentials and Fermi level of the electron in electrolyte 54. In their opinion, the correct description of electrochemical window of electrolyte is not the HOMO-LUMO energy difference, but the potentials of electrolyte reduction and oxidation at the negative and positive potentials, respectively. Nonetheless, the LUMO and HOMO are still important basis to evaluate electrochemical window of electrolyte. Ideally, LIBs operate within this window of LUMO and HOMO. Otherwise, the electrolyte will be reduced on the anode or oxidized on the cathode to form passivating solid electrolyte interphase (SEI) film. Nonetheless, spontaneous side reactions take place on the surface/interface of layered cathodes by contacting with electrolytes 25, 55, 56.

The reaction of the cathode materials with the electrolyte and other inactive cell components

can degrade battery performance, which is likely related to surface structure transformations. Aurbach et al. reported that the existence of nickel element in electrodes augmented the nucleophilicity of the surface oxygen

25.

Electrodes seem to reach passivation in LiPF6/alkyl

carbonate-based solutions, due to the formation of surface species such as LiF, MFx, ROCO2Li, ROCO2M, and polycarbonates.

Figure 6. A schematic view of the Li+ deintercalation from the cathode material. (a) A removal of Li+ is accompanied by the release of an electron from the host material leading to the oxidation of TM. This conditioning is valid when EF of LNMCO lies well above of HOMO of the electrolyte and the downward shift of EF caused by a charging


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potential is still above the HOMO, i.e. no hole transfer to the HOMO takes place. (b) The lowering EF below the HOMO of electrolyte would result in a hole transfer from the TM 3d states to the HOMO until the solid electrolyte interface (SEI) is formed between the cathode and electrolyte. (c) The SEI layer consists of reduced transition metal oxides and the compounds of the decomposed electrolyte. Reprinted with permission from ref 56. Copyright 2015 American Chemical Society.

Cherkashinin et al. reported that high oxidative Ni4+ and Co4+ at high SOC, which are dissolved into electrolyte during redox, are related to the formation of SEI film on cathode material surface 56.

A charging potential leads to the further lowering EF of the cathode material below the HOMO

of the electrolyte, which is accompanied by a hole transfer from TM 3d states to the HOMO of the electrolyte, promoting the decomposition of the oxide and reduction of the TM ions as indicated in Figure 6. As a result, the electronic configurations at the TM sites are changes. And the electrolyte is oxidized to form a SEI-like film (referred as CEI) in the interface, which contains lithium oxides, fluorides, and carbonates, and lowers the electron and lithium ions conductivity. The chemical reactions of NMC materials with the electrolyte also lead to the structural distortion and reduction of the transition metals in the cathode material at the interface impeding the Li+ diffusion from the bulk NMC materials to the electrolyte volume. Thus, the Li+ diffusion rate across the interface can be suppressed or the Li+ paths can even be blocked. The observed reduction of the TM ions at the interface is inherently related to the formation of the CEI layer, which indicates that the highly oxidative Co4+ and Ni3+/4+ ions are involved in CEI layer formation. Additionally, as referred above, the acidic impurity HF produced during hydrolysis of PF6− with trace moisture further corrodes the surface of TCMs. These surface reactions would lead to the formation of a CEI film on the cathode and the dissolution of transition metal species that can poison the anodic SEI. Gao et al. measured the dissolution rates of TMs in 1 M LiPF6/ethylene carbonate (EC) as well as diethyl carbonate (DEC) (1:1 in volume) electrolyte within various cycles by inductively coupled plasma atomic emission spectrometry (ICP-AES) measurement 57. It showed abrupt increases in the dissolutions of TMs, particularly for Ni2+ ions after 100 cycles. Manthiram et al. reported that lithium deposition on graphite anode is triggered by some TMs (especially for manganese) dissolved from the disrupted CEI in the nickel-rich Li[Ni0.61Co0.12Mn0.27]O2/graphite pouch cells 58. The surface structural evolutions triggered by the migration of Ni2+ ions accompanies the creation of hole states at the O2− 2p level, especially the high states-of-charge, which causes oxygen loss at



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the surface of active particles. The released oxygen active-species are so reactive with the carbonatebased electrolyte solvents. Thus, they will produce the thermal runaway which might result in severe safety hazards for the LIBs 59, 60. S Hwang et al. investigated the evolution of the surface structure of LixNi0.8Co0.15Al0.05O2 cathode materials (NCA) as a function of the state of charge at room temperature. There is a greater extraction of lithium from the surface of the particles at over charged state for kinetic effect, resulting in structural instabilities. These instabilities lead to the anomalous reduction of the transition metal ions, the loss of oxygen to maintain charge neutrality, and the formation of new phases at the surface 52. Besides the active materials, the conductive carbon additives seem to react with the carbonate organic electrolyte for various chemical functional groups at the surface. Spontaneous decomposition of electrolyte takes place at the surface region of carbon black particles after storing even without external potential and current

61.

Manthiram et al. reported that organic complexes

generated on carbon during aging migrate to the Li[Ni0.7Co0.15Mn0.15]O2 surface, and to some extent repress these unwanted surface or interfacial reactions during cells working

62.

The passivation

effect by conductive carbon tends to lose stability under high upper cut-off voltages of above 4.5 V, which indicates that the interfacial stability of conductive carbon agents seems to be crucial for cycling stability of advanced high-voltage LIBs. 2.3 Microcracks Microcracks induced by the anisotropic lattice volume change of the layered materials particles, especially the precipitous drop in interlayer spacing at high SOC, could contribute significantly to performance loss

63, 64.

Goonetilleke et al. used operando neutron diffraction to elucidate the

structure evolution of Li(NixCoyMnz)O2 under high voltage 65. Figure 7 shows the evolution of the NMC111 lattice parameters throughout a charge-discharge cycle. The a lattice parameter is observed to expand and contract in an approximately opposite manner to the c lattice parameter. The c-axis increases due to increasing electrostatic repulsion between the transition metal octahedra as the occupancy of the lithium layer is reduced, and the a-axis decreases because of the decrease in average distance between the transition metal atoms caused by transition metal oxidation during charging process until ~4.0 V. In the plateau region (around 4.0 V), the c and a lattice parameter remains constant, which suggests that the average ionic radii of transition metal ions remain the same. Beyond this voltage, the c lattice parameter decrease rapidly due to the migration of transition


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metal ions to lithium layer. Structural collapse under high voltage is observed due to the rapid volume reduction, as shown in Figure 7b, which is caused by the complete extraction of Li ions from between the transition metal layers. Recently, Manthiram reported that the abrupt anisotropic lattice collapse is a universal phenomenon critically dependent on Li utilization, and not Ni content as commonly believed, of layered LiNi1−x−yCoxMnyO2 at deep charge 66. Larger changes in lattice parameter are generally unfavorable as the larger amount of constriction/expansion results in mechanical stress. Therefore, the lattice distortions in opposing directions and the volumetric collapse would generate severe microcracks of material particles, especially at high SOC. Cracks can result in loss of electrical contact between active material particles and the current collector, and expose fresh active material to the electrolyte, which leads to further surface side reactions and loss of active Li ions, formation of surface film with high impedance, and hence, capacity fading upon cycling. Miller et al. revealed that significant separations generated between Li(Ni0.8Co0.15Al0.05)O2 grains even during the very first charge and electrolyte penetration through that crack network into the particle interior (Figure 8ad)

67.

Comparing the results to post-test

microstructural characterization of oxide particles subjected to extensive cycling suggests that the physical separation and isolation of grains may contribute to performance degradation of LIBs. Yan et al. reported unexpected observations on the nucleation and growth of intragranular cracks in the LiNi1/3Mn1/3Co1/3O2 cathode (Figure8 ej)

68.

It is found that the formation of the intragranular

cracks is directly associated with high-voltage cycling, an electrochemically driven and diffusioncontrolled process. Higher cycle voltage will result in deeper Li-ion extraction, which, on one hand, can aggravate structure instability, and on the other hand, can amplify the internal strain within a grain. The intragranular cracks are noticed to be characteristically initiated from the grain interior, a consequence of a dislocation-based crack incubation mechanism. This observation is in sharp contrast with general theoretical models, predicting the initiation of intragranular cracks from grain boundaries or particle surfaces.



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Figure 7. (a) Lattice parameters of NMC 111 as a function of cell voltage during a single charge/discharge cycle. (b) Comparison of change in cell volume as a function of cell voltage for the different NMC compositions. Reprinted with permission from ref 65. Copy right 2018 American Chemical Society.

Figure 8. (a) Cross-section of Li(Ni0.8Co0.15Al0.05)O2 particle in a coin cell after one charge-discharge cycle. Cracking is observed throughout the particle and cracks in several locations are indicated by arrows. (b) Crosssection of Li(Ni0.8Co0.15Al0.05)O2 particle in a coin cell after 4500 charge-discharge cycles showing extensive cracking. (c) Schematic of an idealized particle consisting of several grains indicated by dotted lines. Electrical contact to the shaded conducting medium surrounding the particle is represented by arrows. In (d) after cycling, separations at grain boundaries lead to the loss connectivity and electrical contact between grains. Reprinted with permission from ref 67. Copyright 2013 Wiley. (e–g) charge/discharge profiles of Li/LiNi1/3Mn1/3Co1/3O2 half cells at different high cutoff voltages, and (h–j) low magnification HAADF images of LiNi1/3Mn1/3Co1/3O2 after 100 cycles at different high cutoff voltages. The red arrows indicate voids and the yellow arrows in g indicate intragranular cracks. Scale bars, 500 nm (h–j). Reprinted with permission from ref

68.

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ion+batteries&contentID=10.1038%2Fncomms14101&publicationDate=2017-0116&author=Pengfei+Yan%2C+Jianming+Zheng%2C+Meng+Gu%2C+Jie+Xiao%2C+JiGuang+Zhang+et+al.&imprint=Nature&publisherName=SpringerNature&orderBeanReset=true&volumeNum=8& oa=CC+BY

In summary, the dynamical properties of Li+ transportation is effected enormously by cations mixing, phase transformation and microcrack, which relate to the bulk and surface structure transformation and anisotropic change of lattice parameters under high operating voltage. Meanwhile, the surface reactions are more serious under high operating voltage, resulting in further surface structure degradation. It means that the structure transformation under high voltage is the root of the electrochemical performance deterioration of batteries with TCMs under high operating voltage. Further developments are required to resolve these problems to attain a satisfactory high energy density and long-term cycling stability of the high-voltage TCMs. In the following sections, we will discuss the modification strategies of high-voltage TCMs from the viewpoints of doping, surface coating, synergetic modification of coating and doping, compositional partitioning, electrolyte design and solid battery configuration. 3. MODIFICATIONS Strategies that can modify the intrinsic structural and interface stability will improve the electrochemical performance of TCMs at high voltage. Element doping is the simplest approach for enhancing the intrinsic structural stability by adjusting the crystal lattice at the atomic scale. The surface of cathode material is where electrons or lithium ions diffusion through the materials ends and transfer into conductive agent or electrolyte, thus the majority of the side reactions happen here due to non-equilibrium diffusion reaction 69, 70. These irreversible side reactions are dependent on the species of the cathode surface and the cathode/electrolyte interface environment. So, tuning the surface species of TCMs (coating, compositional partitioning) and interface environment of TCMs/electrolyte (electrolyte design) will affect the irreversible side reactions and thus improve the high-voltage battery performance. 3.1 Doping The mechanisms of improving the high-voltage performance of TCMs from element substitution mainly relate to the decrease of the number of unstable elements, such as Li and Ni; the increase of the bonding energy between oxygen and TMs, resulting in structural stability and decreased oxygen



release; promotion of the lithium ion transport due to increased lithium slab distance by the dopants, and suppression of Ni2+ migration through stabilizing its chemical valence or forming electrostatic repulsion. Besides, sometimes element substitution preferentially replaces Ni and Mn at the surface with more robust cations/anions that suppresses the corrosion of electrolyte and HF under high operating voltage. Therefore, there are many elements have been proposed to doping in TCMs to improve their high-voltage performance, which are summarized in Table 1.


Page 16 of 71

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Table 1. Summary of different dopants and corresponding performance of high-voltage TCMs. (The corresponding anodes are lithium metals) Doping elements

Cation doping

Structure formula

Voltage range

First cycle comparison

Capacity Retention

Current

(V vs. Li/Li+)

(modified vs. pristine)


density

Ref. of

Capacity (mA h g-1)

Rate

Retention (%)

Rate

1C (mA g-1)

Na (1.02Å)

Li0.97Na0.03Ni0.5Co0.2Mn0.3O2

3.04.6

228.43 vs. 189.59

0.1C

-

-

200

71

Cu (0.73Å)

LiNi1/3Co1/3Mn1/3O2

2.54.5

~160 vs. 149

0.5C

75.4 vs. 70, 50cyc

2C

-

72

Mg (0.72Å)

Li[Ni(1/3−z)Co(1/3−z)Mn(1/3−z)Mgz]O2

2.84.4

176 vs. 165

28 mA g-1

98 vs. 95, 30cyc

28 mA g-1

-

73

2.54.5

195 vs. 202

0.1C

67.1 vs. 50.9, 100cyc

0.2C

-

74

(z = 0.04) Ca (1Å)

LiNi0.8(1−x)Co0.1Mn0.1Ca0.8xO2 (x = 6%)

Al (0.535Å)

Li(Ni0.780Co0.100Mn0.068Al0.052)O2

3.04.5

160 vs. 205

20 mA g-1

81 vs. 73, 30cyc

20 mA g-1

-

75

Ti (0.605Å)

Li1.03(Ni0.4Mn0.4Co0.17Ti0.03)O2

2.04.7

~220 vs. 200

0.1 mA/cm2

80 vs. 62.5, 20cyc

0.1 mA/cm2

-

76

Cr (0.62Å)

Li(Ni1/3Co1/3-0.05Mn1/3Cr0.05)O2

2.84.6

~173 vs. 173.3

0.2C

97 vs. 86.6 , 50cyc

0.2C

160

77

Mo (0.62Å)

Li[Ni0.33Mn0.33Co0.33Mo0.01]O2

2.34.6

222 vs. 180

20 mA g-1

83.9 vs. -, 50cyc

20 mA g-1

-

78

Si (0.40Å)

Li[Ni1/3Mn1/3Co1/3]0.96Si0.04O2

2.754.5

180 vs. 174

0.2C

94.7 vs. 85.7, 50cyc

1C

-

79

V (0.59Å)

Li[Ni0.5Co0.2Mn0.3]0.97V0.03O2

2.74.4

222.1 vs. 233.9

0.1C

96.3 vs. 93.3, 50cyc

1C

180

80

Zr (0.84Å)

Li(Ni0.5Co0.2Mn0.3)1-xZrxO2

3.04.6

197.96 vs. 203.55

0.1C

83.78 vs. 69.35,

1C

160

81

100cyc Anion doping

F (1.33Å)

Mixed doping

Fe, Al (0.55Å,

LiNi1/3Co1/3Mn1/3O1.96F0.04

3.04.6

177 vs. 199

0.2C

97.3 vs. 83.5, 30cyc

0.2C

160

82

LiNi0.6Mn0.2Co0.15Al0.025Fe0.02

2.54.4

189 vs. 155

1/20C

61.4 vs.-, 100cyc

1C

-

83

2.84.6

187 vs. 197

20 mA g-1

97 vs. 87, 100cyc

170 mA g-1

-

84

5O2

0.535Å) Mg, F

Li(Ni0.4Co0.2Mn0.4-xMgx)O2-yFy

(0.72Å, 1.33Å )



Page 18 of 71

3.1.1 Single element The commonly used cations for bulk substitution are Na, Cu, Mg, Ca, Al, Ti, Cr, Mo, Si, V, and Zr 71-81, 85-90. The introduction of these cations is suggested to enhance structural stability and reduce oxygen release as well as cation mixing. Na substituted layered Na-NMC material results in a higher specific discharge capacity and superior rate capability with the cut-off voltage of 3.0–4.6 V, which attribute to the incorporation of larger Na+ (rNa+ = 1.02 Å) into the Li+ (rLi+ = 0.76 Å) sites of bulk lattice, enlarging Li layer spacing, decreasing Li+ migration activation energy and reducing the cation mixing 71. Ti4+ (rTi4+ = 0.605 Å) substitution in NMC materials is also helpful in improving structural stability. Doeff et al. reported that the improved high-voltage cyclability (2.0–4.7 V) was achieved when a small percentage of Co was substituted with Ti (4%) (Figure 9a, b)

49, 76.

The

stabilization effects of Ti doping came from the charge compensation by reduction of Mn4+ to Mn3+ in order to balance the valence difference between Co3+ and Ti4+. Meanwhile, Ti substitution can effectively prevent migration of Ni2+ from the transition-metal slab to the Li slab leading to suppression of the phase transition. In addition to the cation substitution of Li or transition metal ions by the doping cations, the anion substitution of O2- by other anions such as F- is also effective to enhance the high voltage performance of TCMs

82, 91-94.

F− with more electronegativity than O2− is broadly utilized in

strengthening the binding energy between the transition metal cations and anions. Furthermore, F− substitution enhances the lattice parameters and it is also pertinent for protecting the surface from HF attack and preventing formation of inhibition raising interphases. It has been reported that a small amount of fluorine-substituted LiNi1/3Co1/3Mn1/3O2-zFz (z = 0.04 and 0.08) exhibit excellent cycling stability and rate capability with the cut-off voltage of 3.0–4.6 V compared to fluorine-free LiNi1/3Co1/3Mn1/3O2 (Figure 9c, d)

82.

Consequently, the O2− substitution by F− significantly

enhances the electrochemical properties like capacity retention, rate capability, as well as thermal stability.


Page 19 of 71 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 9. Capacity as a function of cycle number for lithium cells containing selected electrodes from the (a) NMC333-GNC (glycine nitrate combustion), COH (co-precipitation), and nano Ti series and (b) NMC442-GNC and COH series, charged and discharged at 0.1 mA/cm2 between 2.0–4.7 V. Reprinted with permission from ref

76.

Copyright 2012 The Electrochemical Society. (c) Initial charge-discharge curves of LiNi1/3Co1/3Mn1/3O2-zFz samples in the voltage range 3–4.6 V at a charge/discharge current density of 32 mA g-1 (0.2 C) at room temperature. (d) Cycling performance of LiNi1/3Co1/3Mn1/3O2-zFz samples between 3 and 4.6 V at 0.2 C rate at room temperature. Reprinted with permission from ref 82. Copyright 2007 Elsevier.

3.1.2 Mixed elements The doped inactive cations do not participate in the redox process, which maintain the interslab space during Li+ de/intercalation and improve the structural stability of TCMs. The anion doping can protect the surface from HF attack. Therefore, the mixed elements doping can combine their advantages to improve the stability of intrinsic structure and interface. Yang-Kook Sun et al. reported a material of Li[Ni0.4Co0.2Mn(0.4−x)Mgx]O2−yFy (x = 0.04, y = 0.08) with fluorine substitution for oxygen and magnesium for manganese 84. Substituting Mg for part of the metal in the transition metal layer improves the structural stability and F substitution in the oxygen sites protects the cathode particle surface from HF attack. Therefore, the co-doping of Mg and F would provide a



synergistic

effect

in

improving

the

Page 20 of 71

electrochemical

performance.

Although

the

Li[Ni0.4Co0.2Mn0.36Mg0.04]O1.92F0.08 delivered somewhat slightly lower initial discharge capacity the capacity retention, interfacial resistance, and thermal stability were greatly enhanced comparing to the Li[Ni0.4Co0.2Mn0.4]O2 and Li[Ni0.4Co0.2Mn0.36Mg0.04]O2 (Figure 10).

Figure 10. (a) The discharge capacity vs. number of cycles of: a Li/Li[Ni0.4Co0.2Mn0.4]O2, b Li/Li[Ni0.4Co0.2Mn0.36Mg0.04]O2, and c Li/Li[Ni0.4Co0.2Mn0.36 Mg0.04]O1.92F0.08 cells in the voltage range of 2.8–4.6 V.

(b)

Differential

scanning

calorimetry

(DSC)

profiles

of:

a

Li/Li[Ni0.4Co0.2Mn0.4]O2,

b

Li/Li[Ni0.4Co0.2Mn0.36Mg0.04]O2, and c Li/Li[Ni0.4Co0.2Mn0.36Mg0.04]O1.92F0.08 cells charged to 4.3V. Nyquist plots of: (c) the Li/Li[Ni0.4Co0.2Mn0.4]O2 and (d) Li/Li[Ni0.4Co0.2Mn0.36Mg0.04]O1.92F0.08 at different cycles of the 5th, 10th, 15th, and 20th. Reprinted with permission from ref 84. Copyright 2006 Elsevier.

Element doping is an effective and simple method to improve the structural stability by mitigating the collapse of the interlayer spacing upon high delithiation, and the charge compensation of substitutions may increase the capacity. However, inactive element doping sometimes involves the capacity sacrifice when the doping element content increases. Therefore the doping content needs to be controlled by a suitable amount to maintain the balance of capacity and cycling stability. Although many elements have been reported as dopants for TCMs, how the electrochemical


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properties change with the dopants clearly need to be verified. In addition, from a commercialization perspective, cost of the dopants and doping methods should be considered. Therefore, more fundamental studies of the doping effect and economical doping method should be conducted to further develop the high-voltage TCMs. 3.2 Surface coating As mentioned above, the side reactions between the active material and electrolyte have significant effect on the performance of the cathodes. To protect the electrode from directly contacting organic electrolyte, coating layers are commonly applied. The main purpose of the surface coating is to reduce potential side reactions and their effects, including transformations to rocksalt/spinel, the dissolution of transition-metal ions, the oxygen release, and the formation of an undesirable SEI-like film on the cathodes surface at high working voltage. Many coating materials have been reported to improve the high-voltage performance of TCMs, such as inactive coating including oxide, fluorides, and phosphates, electron conductive coating including carbon-based materials, conducting polymers, and lithium ion conductor coating. A summary of various coating materials and their effect on the electrochemical performance of TCMs with high operating voltage are shown in Table 2.



Page 22 of 71

Table 2. Summary of different surface coatings and corresponding high-voltage performance of TCMs. (The corresponding anodes are lithium metals.) Coating materials

Inactive coatings

Structure formula

Oxides

Fluorides

Voltage range


Capacity Retention

(V vs. Li/Li+)



Current density of

Ref.

1C rate (mA g-1)

Capacity (mA h g-1)

Rate

Retention (%)

Rate

Al2O3

LiNi0.5Co0.2Mn0.3O2

3.04.5

162 vs. 158

0.5C

85 vs. 75, 100cyc

0.5C

160

95

CeO2

LiNi0.5Co0.2Mn0.3O2

2.84.6

234.5 vs. 183.8

0.05C

57.7 vs. 15.4, 100cyc

10C

180

96

CuO

LiNi0.5Co0.2Mn0.3O2

3.04.8

160 vs.150

1C

43 vs. 39, 50cyc

1C

190

97

Sb2O3

LiNi0.4Co0.2Mn0.4O2

3.04.6

177 vs. 185

0.5C

93.8 vs. 73.3, 100cyc

0.5C

200

98

TiO2

LiNi0.6Co0.2Mn0.2O2

3.04.5

193.9 vs. 187.6

0.2C

88.7 vs. 77.1, 50cyc

1C

140

99

ZnO

LiNi0.5Co0.2Mn0.3O2

2.54.5

238.1 vs. 201.2

0.1C

91.5 vs. 87.4, 60cyc

2C

200

100

ZrO2

LiNi1/3Co1/3Mn1/3O2

3.0-4.5

173.0 vs. 173.5

0.2C

99.1 vs. 78.8, 100cyc

0.2C

150

101

SiO2

LiNi0.5Co0.2Mn0.3O2

2.84.5

165 vs. 172

1C

93.4 vs. 66.0, 100cyc

1C

-

102

Bi2O3

LiNi0.4Co0.2Mn0.4O2

2.54.6

214 vs. 226

0.1C

81.7 vs.65.9, 100cyc

0.1C

-

103

In2O3

LiNi0.4Co0.2Mn0.4O2

2.54.6

221 vs. 226

0.1C

68.7 vs.65.9, 100cyc

0.1C

-

103

Cr2O3

LiNi1/3Co1/3Mn1/3O2

2.84.5

196.3 vs. 196.8

0.1C

83 vs. 72.4, 30cyc

0.5C

200

104

Y2O3

LiNi0.5Co0.2Mn0.3O2

3.04.6

182.5 vs. 185.9

1C

87.8 vs 78.6, 50cyc

1C

180

105

Nb2O5

LiNi1/3Co1/3Mn1/3O2

2.84.6

~160 vs. 155

1C

64 vs. 0, 100cyc

1C

2.31

106

AlF3

LiNi1/3Co1/3Mn1/3O2

3.04.5

172 vs. 173

0.5C

93 vs. 75, 50cyc

0.5C

160

107

(NH4)3AlF6

LiNi1/3Co1/3Mn1/3O2

3.04.5

192 vs. 192

0.2C

92 vs. 76, 70cyc

0.5C

160

108

ZrFx

LiNi1/3Co1/3Mn1/3O2

3.04.6

176.96 vs. 170.52

1C

-

-

-

109

LaF3

LiNi0.5Co0.2Mn0.3O2

3.04.8

~220 vs. 210

1C

~77.2 vs. 75.2, 50cyc

1C

200

110

MgF2

LiNi0.5Co0.2Mn0.3O2

3.04.8

~200 vs. 210

1C

~90 vs. 75.2, 50cyc

1C

200

110

FeF3

LiNi1/3Co1/3Mn1/3O2

3.04.8

176 vs. 151

200 mA g-1

84 vs. 57, 50cyc

200 mA g-1

-

111

CaF2

LiNi1/3Co1/3Mn1/3O2

2.54.5

187.1 vs. 191.0

0.1C

98.1 vs. 83.3, 50cyc

0.1C

280

112

SrF2

LiNi1/3Co1/3Mn1/3O2

2.54.6

165.7 vs.179

-

86.9 vs. 76.3, 50cyc

-

-

113

YF3

LiNi1/3Co1/3Mn1/3O2

3.04.5

112 vs. 121

5C

93 vs. 74, 100cyc

5C

200

114

FePO4

LiNi0.5Co0.2Mn0.3O2

2.54.6

213 vs. 214

0.2C

91.2 vs. 82.1, 50cyc

1C

-

115

MnPO4

LiNi0.6Co0.2Mn0.2O2

3.04.6

221.0 vs 215.8

0.1C

86.9 vs. 67.4, 50cyc

0.1C

-

116

(55C)

Phosphates


Page 23 of 71 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Electron

Carbon-based

C

LiNi1/3Co1/3Mn1/3O2

2.54.5

191.2 vs. 175.7

0.5C

90.3 vs. 72.4, 100cyc

1C

170

117

conductive coatings

materials

NC

LiNi0.6Co0.2Mn0.2O2

3.04.5

192.1 vs. 202.4

0.2C

92.0 vs.76.4, 100cyc

1C

195

118

3.04.5

183.4 vs. 181.9

1C

85.6 vs. 60.4, 100cyc

1C

-

119

rGO

LiNi0.6Co0.2Mn0.2O2

PPy

LiNi0.8Co0.1Mn0.1O2

2.84.5

179 vs. 172

2C

66 vs. 58, 100cyc

2C

18

120

polymers

PEDOT

LiNi1/3Co1/3Mn1/3O2

2.84.5

156 vs. 168

150 mA g-1

85.9 vs. 69.6, 80cyc

150 mA g-1

-

121

Inorganic conductors

LiF

LiNi1/3Co1/3Mn1/3O2

2.54.5

187.2 vs.189.8

0.1C

78.7 vs. 44.9, 50cyc

10C

280

122

LiTaO3

LiNi1/3Co1/3Mn1/3O2

3.04.8

167 vs. 182

160 mA g-1

73 vs. 32, 100cyc

160 mA g-1

-

123

Li2ZrO3

LiNi0.7Co0.15Mn0.15O2

3.04.5

190 vs. 190

1/3C

92.8 vs. 82.4, 100cyc

1/3C

180

124

Li2SiO3

LiNi0.5Co0.2Mn0.3O2

3.04.5

190.1 vs. 195.9

0.2C

96.1 vs. 55.3, 100cyc

10C

-

125

Li2SnO3

LiNi0.5Co0.2Mn0.3O2

3.04.5

134 vs.120

10C

94.9 vs 68.3, 100cyc

10C

180

126

Li2TiO3

LiNi0.5Co0.2Mn0.3O2

3.04.6

182.2 vs. 187.8

0.2C

92.4 vs. 74.3, 100cyc

1C

140

127

Li3VO4

LiNi0.5Co0.2Mn0.3O2

2.84.6

148.9 vs. 178.5

10C

41.3 vs. 1.4, 100cyc

10C

180

128

LiAlO2

LiNi1/3Co1/3Mn1/3O2

2.84.5

177.2 vs. 173.9

0.2C

96.7 vs. 82.3, 50cyc

1C

-

129


2.54.5

147 vs. 118

200 mA g-1

89 vs. 64, 50cyc

200 mA g-1

-

130

LiNi0.4Co0.3Mn0.3O2

3.04.8

190.1 vs. 196

1C

74.3 vs. 64.1, 50cyc

1C

-

131


2.74.5

~226 vs. 221

C/5

91.6 vs. 79.0, 200cyc

1/3C

200

132

NaTi2(PO4)3

LiNi0.6Co0.2Mn0.2O2

3.04.6

214.2 vs. 208.8

0.2C

85.3 vs. 49.0, 100cyc

1/2C

180

133

LiBO2

LiNi0.5Co0.2Mn0.3O2

3.04.5

184.38 vs. 191.37

0.2C

91.23 vs. 53.19, 50cyc

1C

160

134

LiFePO4

LiNi0.5Co0.2Mn0.3O2

3.04.5

181.5 vs. 189.7

1/3C

92.8 vs. 81.9, 150cyc

1/3C

-

135

Li3Fe2(PO4)3

LiNi0.5Co0.2Mn0.3O2

2.84.5

186.0 vs. 189.1

0.1C

75.33 vs. 51.92, 100cyc

1C

-

136

PI

LiNi1/3Co1/3Mn1/3O2

2.84.6

~180 vs 180

0.5C

71 vs. 58, 50cyc

1C

2.36 mA cm-1

137

Li-ion and electron

LixAlO2&LixTi2O4

LiNi0.8Co0.1Mn0.1O2

3.04.4

213 vs. 208

0.2C

98 vs. 80, 100cyc

0.5C

-

138

conductive coatings

Li3PO4&PPy

LiNi0.8Co0.1Mn0.1O2

2.84.5

203 vs. 206

0.1C

95.1 vs. 86, 50cyc

0.1C

200

139

Li3VO4&PPy

LiNi0.6Co0.2Mn0.2O2

2.74.5

185.6 vs. 184.8

0.5C

93.7 vs.73.6, 100cyc

0.5C

180

140

Electron

Li-ion conductive

conductive

coatings

LiLaTiO3 Li3PO4 Li3PO4(infused)

Li-ion conductive polymers



Page 24 of 71

3.2.1 Inactive coating Coating of electrochemically and chemically inactive material on the cathode surface can act as a physical protection barrier to suppress the parasitic reactions and hence improve the electrochemical performance of LIBs. The common coating materials mainly contain oxides 95-106, 141-146(Al

2O3,

fluorides 116(FePO

CeO2, CuO, Sb2O3, TiO2, ZnO, ZrO2, SiO2, Bi2O3, In2O3, Cr2O3, Y2O3, Nb2O5),

107-114(AlF 4,

3,

(NH4)3AlF6, ZrFx, LaF3, MgF2, FeF3, CaF2, SrF2, YF3), phosphates

115,

MnPO4). Moreover, the modification matrix can act as an effective HF scavenger and

neutralize the acidity near the electrode/electrolyte interface. The oxides can react with HF to reduce HF concentration in the electrolyte, such as Al2O3+6HF→2AlF3+3H2O. The by-product of water facilitates HF generation in electrolyte, and further corrodes electrode surface. Fortunately, another by-product of metal fluoride is very stable and can protect the cathode materials from further corrosion following complete consumption of the metal oxides. Compared to the oxides, the fluorides and phosphates coating are more stable owing to their intrinsic nature. Su Hyun Yun et al. reported that the ZrFx-coated Li[Ni1/3Co1/3Mn1/3]O2 have a better thermal stability, and the discharge capacities and rate capabilities of coated Li[Ni1/3Co1/3Mn1/3]O2 electrodes in the voltage range of 3.04.6 V have been improved (Figure 11a, b) 109. Kim et al. introduced FeF3 as a coating material for the surface modification of the Li[Ni1/3Co1/3Mn1/3]O2 cathode. The sample coated with FeF3 exhibited a higher discharge capacity and improved rate capability than the pristine sample. Additionally, the cyclic performance at a high voltage range (3.04.8 V) was also enhanced by the FeF3 coating (Figure 11c, d) 111. Yao et al. have developed a phosphate surface modification route via H3PO4 etching to improve the electrochemical performances of porous LiNi1/3Co1/3Mn1/3O2 microspheres synthesized by a co-precipitation approach combined with high-temperature calcinations

147.

The enhanced rate capability and cycle stability of phosphate-coated

LiNi1/3Co1/3Mn1/3O2 microspheres could be attributed to the phosphate surface modification, which suppressed the side reaction between cathode and electrolyte (Figure 11e, f).


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Figure 11. (a) Initial discharge profiles of pristine and ZrFx-coated Li[Ni1/3Co1/3Mn1/3]O2 electrodes in the voltage range of 3.04.6 at 6 C. (b) Discharge capacities and cyclic performances of pristine and coated Li[Ni1/3Co1/3Mn1/3]O2 electrodes in the voltage range of 3.04.6 at 1, 2, 4, 6, 8, and 12 C rates. Reprinted with permission from ref 109. Copyright 2010 Elsevier. (c) Discharge capacities and cyclic performances of pristine and coated Li [Ni1/3Co1/3Mn1/3]O2 electrodes in the voltage range of 4.63.0 V at current densities of 40, 100, 200, 400, and 600 mA g-1. (d) Cyclic performances of the pristine and FeF3-coated Li[Ni1/3Co1/3Mn1/3]O2 electrodes at a current density of 200 mA g-1 in the voltage range of 4.83.0 V at room temperature. Reprinted with permission from ref 111.

Copyright 2014 Elsevier. (e) The rate capability of the pristine and phosphate-coated LiNi1/3Co1/3Mn1/3O2

microspheres at different current rates. (f) Nyquist plots of the pristine and 2 mol.% H3PO4 modified LiNi1/3Co1/3Mn1/3O2 microspheres and equivalent circuit model. Reprinted with permission from ref 147. Copyright 2013 Elsevier.

Inactive coating can suppress the corrosion of electrolyte effectively, alleviating surface phase transition of TCMs particles, thus improving the cycling performance. But the rate capability of these coated materials is not as obviously improved as the cycling performance for the insulating nature. Besides, the thickness of an inactive coating is a critical factor because these coating materials are electrochemically inactive and electronically insulating. Coating layers need to remain relatively thin to enable facile Li-ion transport, otherwise it would lower a cycling capacity and rate performance for both low electronic and lithium ions conductivity. When such a layer is too thin, it cannot provide a sufficient protection against metal dissolution and parasitic reactions of TCMs particles and electrolytes. Therefore, the electron or/and lithium ion conductive materials would be



Page 26 of 71

better choices for the surface coating. 3.2.2 Electron, lithium ion and dual-conductive conductor coating Considering the electronically insulating of inactive coating materials with poor rate performance, the electron conducting materials are proposed as coating layers. The electron conducting coating layer can improve the electron transfer between TCMs particles, accelerating the charge transfer process on the NMC surface. The electron conducting layers largely contain carbon, graphene oxide and conducting polymers

117-121, 140, 148, 149.

Xiong et al. reported that the conducting polypyrrole

(PPy) has been successfully coated on the surface of LiNi0.8Co0.1Mn0.1O2 through a new and facile approach

120.

When cycling at high temperature and high voltage, PPy-coated material exhibited

noticeably improved performance compared with that of pristine material (Figure 12a, b). Such improvements are ascribed to the conducting PPy coating layer, not only facilitating the migration of electron, but also alleviating the reaction between the highly delithiated Li1−δNi0.8Co0.1Mn0.1O2 and the electrolyte at a high voltage. Lithium-ion conductors are also broadly applied as coating materials for cathode surfaces, which mainly contain LiF, LiTaO3, Li2ZrO3, Li2SiO3, Li2SnO3, Li2TiO3, Li3VO4, LiAlO2, LiLaTiO3, Li3PO4, NaTi2(PO4)3, LiBO2, LiFeO4 and Li3Fe2(PO4)3

122-136, 150.

Li et al. reported a uniform and

dense lithium conductive LiTaO3 layer with controllable layer thicknesses coated on the NMC by ALD. The inactive coated materials often demonstrate reduced ionic conductivity, as well as hindered lithium ion transport into and out of the electrode particles, resulting in limited performance improvement. By contrast, a solid-state electrolyte coating may overcome the limited lithium ion conductivity (Figure 12c). It indicates that the LiTaO3 coating is effective to improve cycling performance of the cathode with a cutoff potential of 3.04.8 V (Figure 12d) 123. Song et al. reported a long-life layered oxide cathode (LiNi0.7Co0.15Mn0.15O2) with a uniform surface coating of the cathode particles with Li2ZrO3. A pouch-type full cell fabricated with the Li2ZrO3-coated cathode and a graphite anode displays 73.3% capacity retention after 1500 cycles at a C/3 rate with the cut-off voltage of 2.54.4 V (Figure 12e, f) 124.


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Figure 12. (a) The discharge capacities of the pristine and PPy-coated NMC at various currents (0.110 C). Inset is the initial discharge curves at different C rates. (b) The cycling performance of pristine and PPy-coated sample between 2.8 and 4.5 V at 25 °C. Insets indicate the discharge profiles at different cycles. Reprinted with permission from ref 120. Copyright 2014 Springer Nature. (c) The difference between solid state electrolyte coatings and metal oxide coatings. The former provides higher lithium ion conductivity than the latter. (d) The comparison of cyclic performance of NMC-0, NMC-2, NMC-5, NMC-10, and NMC-20 at a current density of 160 mA g-1 at room temperature in the voltage range of 3.04.8 V (NMC cathode electrodes were coated by 0, 2, 5, 10 and 20-ALD cycle LiTaO3). Reprinted with permission from ref 123. Copyright 2014 The Royal Society of Chemistry. (e) Longterm cycle performance of a pouch-type full cell assembled with graphite as an anode and (f) its corresponding charge/discharge profiles. The testing was performed between 2.5 and 4.4 V at a C/3 rate and 25 °C. Reprinted with permission from ref 124. Copy right 2017 American Chemical Society.



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Apart from these inorganic metal oxides, some ionic conductive polymers are also used as surface coating materials. The inorganic coating materials are brittle and weak adhesion while polymers with suited mechanical properties may alleviate the delamination of coating layer. Jang-Hoon Park et al. have reported polyimide (PI) as the surface coating materials to improve the electrochemical performance of LiNi1/3Co1/3Mn1/3O2 137. With the PI coating layer, the capacity retention ratios at higher cut-off voltages LiNi1/3Co1/3Mn1/3O2 increased from 58% to 71% (2.84.6 V) and 52% to 66% (2.84.8 V) after 50 cycles at 1 C, respectively. The enhanced electrochemical performance was attributed to the presence of PI layer which can effectively prevent the formation of undesired resistive layer from accelerating the resistance growth (Figure 13). Meanwhile, the thermal stability of PI coated LiNi1/3Co1/3Mn1/3O2 was significantly enhanced due to the stable thermal and chemical properties of PI.

Figure 13. Discharge capacities (charge/discharge current density = 1.0 C/1.0 C) as a function of cycle number for cells assembled with pristine or PI-wrapped LiNi1/3Co1/3Mn1/3O2 under a voltage range between: (a) 2.8 and 4.6 V; (b) 2.8 and 4.8 V. The insets indicate the discharge profiles of cells after the 1st and 50th cycles at a given voltage range. Variation in AC impedance spectra (1st/50th cycles) of cells: (c) Pristine LiNi1/3Co1/3Mn1/3O2 (2.84.6 V); (d) Pristine LiNi1/3Co1/3Mn1/3O2 (2.84.8 V); (e) PI-wrapped LiNi1/3Co1/3Mn1/3O2 (2.84.6 V); (f) PI-wrapped LiNi1/3Co1/3Mn1/3O2 (2.84.8 V). Reprinted with permission from ref

137.

Copy right 2012 The Royal Society of

Chemistry.

The dual-conductive coating layers can afford high ionic and electronic conductivity 139, 140. Wu et al. combined Li3PO4 with PPy to form dual conductive coating layers on LiNi0.8Co0.1Mn0.1O2


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cathode material for improving the cycling and rate performance

139.

The Li3PO4 layer would

remarkably improve the ionic conductivity of the cathode materials and reduce the generation of HF. The PPy layer could form a uniform film which can make up for the Li3PO4 coating defects and enhance the electronic conductivity. The stretchy PPy capsule shell can reduce the generation of internal cracks by resisting the internal pressure as well. Thus, ionic and electronic conductivity, as well as surface structure stability have been enhanced after the modification. The capacity retention of the modified cathode material is 95.1% at 0.1 C after 50 cycles between 2.84.5 V, whereas the bare sample is only 86%, and performs 159.7 mA h g-1 at 10 C compared with 125.7 mA h g-1 for the bare. Besides the surface coating, the “intergranular coating” seems also important for the electrochemical performance of TCMs. Zhang et al. infused a solid electrolyte, Li3PO4 into the grain boundaries of LiNi0.76Mn0.14Co0.10O2 to enhance the capacity retention and voltage stability 132. The solid electrolyte infused in the boundaries not only acts as a fast channel for lithium ion transport but also prevents penetration of the liquid electrolyte into the boundaries, and consequently eliminates the detrimental factors, which include cathode-liquid electrolyte interfacial reactions, intergranular cracking and layered-to-spinel phase transformation. The Li3PO4-infused electrode demonstrated the highest capacity retention of 91.6% after 200 cycles within a voltage window of about 2.74.5 V, in contrast to 79.0% for pristine particles (Figure 14).



Figure 14. (a) The schematic illustration shows the evolution of the Li3PO4 (LPO) coating layer on the secondary particle following coating and annealing. (b) Cycling performance of the three Ni-rich NMC cathodes in the voltage range 2.74.5 V at room temperature. The half-cell is cycled at C/3 after three formation cycles at C/10. Inset: the magnified area (red frame) shows that the capacity retention is 91.6% for the LPO-infused cathode and 79.0% for the pristine cathode. Cross-sectional SEM image of (c) the pristine electrode and (d) the LPO-infused electrode after 200 cycles. (e) Impedance spectra evolution of the pristine electrode and the LPO-infused electrode at the (e) 1st and (f) 50th cycles. Reprinted with permission from ref 132. Copy right 2018 Springer Nature.

Though electron and lithium ion conductive coatings can overcome the disadvantages of inactive coatings due to their conductive properties, their uniformity and thickness also influence the capacity and performance of TCMs. The suitable method to control the thickness of coatings is the thin film technology such as chemical vapor deposition (CVD), atomic layer deposition (ALD), sputtering. However, it increases the cost of the production of these cathode materials using in the power batteries, which is a critical factor limiting their practical using. In addition, the surface coating can improve the surface properties that result from the side reaction of electrolyte and TCMs particles, but it is supposed to be not as effective in suppressing the phase transition that occur upon electrochemically cycling, such as cations mixing and O2 release. 3.3 Synergetic modification of coating and doping


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As we discussed above, the interior structural stability and interface stability of TCMs are equally important for their electrochemical performance. Surface coating and element doping are effective means to enhance the performance of TCMs. However, the individual modification can only solve parts of the issues. Hence, the combination of coating and elemental doping has drawn more attention of researchers. The synergetic modification of coating and doping would combine their advantages to increase the structural stability and suppress the interface reaction and HF corrosion at high voltage, simultaneously. Some synergetic modifications and their effects on electrochemical performance of TCMs are summarized in table 3.



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Table 3. Summary of the synergetic modifications of coating and doping and corresponding performance of high-voltage TCMs. (The corresponding anodes are lithium metals.)

Modification

Structure formula

Voltage range


Capacity Retention

(V vs. Li/Li+)



Current density of

Ref.

1C (mA g-1)

Capacity (mA h g-1)

Rate

Retention (%)

Rate

F− doping and LiF coating

LiNi0.5Co0.2Mn0.3O2

3.04.5

191.4 vs. 187

0.5C

81.1 vs. 35.1, 300cyc

5C

150

151

Zr4+/Ti4+

LiNi0.5Co0.2Mn0.3O2

3.04.4

202.1 vs. 203.8

0.1C

94.20 vs. 75.59, 200cyc

1C

155

152

LiNi0.6Co0.2Mn0.2O2

2.54.5

156.8 vs. 141.2

0.5C

79.85 vs. 77.56, 100cyc

0.5C

160

153

LiNi0.8Co0.1Mn0.1O2

2.84.5

190 vs. 207.4

0.5C

98.4 vs. 84.5, 100cyc

0.5C

200

154

LiNi0.5Co0.2Mn0.3O2

3.04.6

186.6 vs. 186.6

1C

87.4 vs. 74.1, 100cyc

1C

155

155

LiNi1/3Co1/3Mn1/3O2

2.54.5

173.9 vs. 172.0

1C

79.0 vs. 61.7, 300cyc

1C

200

156

LiNi0.6Co0.2Mn0.2O2

3.04.6

200.8 vs. 199.9

1C

82.63 vs. 66.03, 100cyc

1C

165

157

LiNi0.8Co0.1Mn0.1O2

2.754.5

196.8 vs. 178.7

0.2C

86.9 vs. 70.7, 100cyc

5C

200

158

LiNi0.8Co0.1Mn0.1O2

2.84.5

~208 vs. 209.4

0.1C

83.2 vs. 69.9, 200cyc

1C

200

159

doping

and

Li4Ti5O12/Li2TiO3 coating Si4+ doping and Li2SiO3 coating Y3+ doping and LiYO2 coating Al3+ doping and Al2O3 coating La3+ doping and La2O3 coating Cd2+ doping and CdO coating Ce4+ doping and CeO2 coating Zr4+ doping and Li2ZrO3 coating


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Liu et al. reported that both surface F− doping and LiF coating, with LiPF6 as one precursor, were successfully carried out at low calcination temperature in air to modify the surface of LiNi0.5Co0.2Mn0.3O2

151.

Surface F− doping and LiF coating simultaneously triggers partial Ni3+

reduction to Ni2+ and partial M−O bond is replaced by a higher-energy M−F bond. Meanwhile, F− doping enlarges the specific surface area of LiNi0.5Co0.2Mn0.3O2 and LiF coating protects materials from side reactions with electrolytes, which enhance the cyclic stability and rate capacity. The modified material retains 93.7% of its initial capacity and delivers 179.4 mA h g-1 at a current density of 0.5 C after 100 cycles at 3.04.5 V (Figure 15a, b). Chen et al. reported the synergetic advantages of doping and coating by zirconium and titanium elements 152. The Zr atoms have been introduced into the LiNi0.5Co0.2Mn0.3O2 cathode material, which stabilizes its host structure. Part of the introduced Ti not only dopes into the Zr-doped LiNi0.5Co0.2Mn0.3O2 cathode material, and the other Ti reacts with the lithium residues on the surface of LiNi0.5Co0.2Mn0.3O2 cathode material to form the Li4Ti5O12/Li2TiO3 mixed coating layers, which significantly improve its interfacial structure and kinetic characteristics. The capacity retentions of the Zr/Ti co-modified LiNi0.5Co0.2Mn0.3O2 reach to 94.20% after 200 cycles over 3.04.4 V and 91.71% after 100 cycles over 3.04.6 V (Figure 15c, d).

Figure 15. (a) Schematic illustrations of the fabrication of modified LNCM. (b) Cycle performance of LNCM, LNCM-1, LNCM-2, and LNCM-3 at 0.5 C at 3−4.5 V vs Li/Li+. The content of LiPF6 in the precursor mixture was 0, 0.5, 0.7, and 1 wt %, and attained products were denoted as LNCM, LNCM-1, LNCM-2, and LNCM-3, respectively. Reprinted with permission from ref

151.

Copyright 2018 American Chemical Society. (c) Schematic

illustration of the structural/interfacial evolution of the LiNi0.5Co0.2Mn0.3O2 after Zr and Ti co-modification; and (d) Cycle performance of the as-prepared samples at the rate of 1C over 3.04.6 V. Reprinted with permission from ref



152.

Copyright 2018 Elsevier.

Chen et al. reported a Cd-modified LiNi0.6Co0.2Mn0.2O2 cathode material with Cd doping and CdO coating 157. Unlike the doping approach we mentioned above, the Cd diffused into the crystal lattice of the LiNi0.6Co0.2Mn0.2O2 with gradient contribution. The Cd gradient doping stabilizes the crystal structure from surface to bulk and expands the transportation tunnels of electron and Li+ (Figure 16a). And the cadmium oxide coating layers improve the interfacial properties, which is favorable to effectively avoid the active Ni4+ reacting with the electrolyte and the impedance increasing at high voltage upon cycling. As a result, the better cycling stability and rate capability of the cathode at high voltage after cadmium oxide modification are obtained (Figure 16b, c). Wu et al. reported an approach to improve the electrochemical performances of LiNi0.8Co0.1Mn0.1O2 microspheres by gradient doped Zr4+ 160. The doped Zr4+ presented a gradient distribution from the surface to the bulk of Li1-xZrx(Ni0.8Co0.1Mn0.1)1-yZryO2 and occupied both the transition-metal slab and Li slab. The gradient doping can take advantage of the “pillar effect” while restraining the “blocking effect” to the maximum extent, thus reduce irreversible capacity loss and improve the cycling and rate performance of the Ni-rich cathode materials. Moreover, when the doping amount is increased, some Zr4+ can form Li2ZrO3 on the surface as a coating layer to enhance the ionic conductivity of the cathode and inhibit side reactions between the cathode surface and the electrolyte, which increases the cycle stability and rate performance of the Ni-rich cathode materials (Figure 16 d). The capacity retention of the modified sample reached 83.2% after 200 cycles at 1 C (200 mA g-1) at 2.8–4.5 V, and the discharge capacity was up to 164.7 mA h g-1 at 10 C (Figure 16e, f). This effective strategy shows us a new method with gradient doping to improve the NMC performance. However, the degree of structure stabilization brought about by the concentration gradient is still unknown.


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Figure 16. (a) Schematic illustrations of the microstructure evolution and the reaction mechanisms of pristine and Cd-modified LiNi0.6Co0.2Mn0.2O2. (b) Cycle performance of the NCMs at the rate of 1 C at 3.0–4.6 V and 1C &8C at 3.0–4.5 V. Reprinted with permission from ref xZrx(Ni0.8Co0.1Mn0.1)1-yZryO2

157.

Copyright 2018 Elsevier. (c) Schematic diagram of Li1-

(0≤x+y≤0.02). (d) Cycle performance and rate properties for Zr-0, Zr-0.005, Zr-0.01,

and Zr-0.02 at 2.8–4.5 V. Reprinted with permission from ref 159. Copyright 2018 Wiley.

3.4 Compositional Partitioning As mentioned above, the different elements in TCMs have different influence on the performance of LIBs. Mn and Al are effective in suppressing undesired surface chemical and structural evolution of the material at high voltage, and Co is commonly incorporated to improve layered ordering and electrical conductivity. The high content of Ni would enhance the capacity but introduce cation mixing and phase transformation when the battery under high operating voltage. Therefore, in order to obtain both high capacity and high structural stability, the design of chemical compositional variation from the interior to the exterior of a micron-sized secondary particle was developed. Compositionally partitioning of the TCMs usually places a Ni-rich composition in the particle center and a Ni-deficient composition at the particle surface, which provides a high capacity and high stable interface. This has been done with different configurations, including core−shell 161-166, core −shell with concentration gradient 165, 167, and full concentration gradient (FCG) 168-172. Core–shell structure is regarded as the double or multiple layers where the interior component (core) and the outer surface (shell) have different chemical components, thus enabling the recombination and complementation of the benefits of core and shell compounds. The core materials



are always Ni-rich materials and the shell layers are cathode materials enriched with Mn, which have high energy density and good structural and thermal stability, respectively. This will significantly improve the structural stability during cycling by partially reducing the Ni3+ ions to more thermodynamically stable Ni2+ 173. On the basis of this idea, a fully functional microscale core/shell cathode was realized using a coprecipitation method. Yonghyun Cho et al. reported a high-capacity and safe cathode material with an average composition of Li[Ni0.54Co0.12Mn0.34]O2, in which the shell and core consist of the highly stable spinel and the layered phase, respectively (Figure 17)

164.

The material demonstrated a reversible capacity of 200 mA h g−1 with the cut-off

voltage of 3.04.5 V and retained 95% capacity retention at the most severe test conditions of 60 °C. In addition, the amount of oxygen evolution from the lattice in the cathode with the core-shell was reduced by 70%. This outstanding performance is attributed to the rational integration of the advantageous features of the structural stability of the core and the chemical stability of the shell, effectively addressing impedance rise and capacity fade stemming from the aggressive reaction between the cathode surface and the organic electrolyte at the higher operating voltages. Core–shell structures could be considered a conceptual extension of surface coatings. The coating layer is replaced with a Mn-rich cathode material which is electrochemical active and possesses better thermal and structural stability. Unlike the surface coating, the shell thickness is not controlled strictly because the shell is electrochemical active and lithium and electronic conductive.

Figure 17. (a) Schematic view of a core-shell particle with heterostructures. (b) Schematic diagrams for the preparation of the core-shells (HS-LiNi0.54Co0.12Mn0.34O2). c) SEM images of the HS-LiNi0.54Co0.12Mn0.34O2, corresponding to (b). Reprinted with permission from ref 164. Copyright 2011 Wiley.

Though the core-shell structure improves the performance of the layered cathode materials, the structural mismatch between the core and the shell can lead to void formation at the core/shell


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interface after long-term cycling under high voltage, blocking Li+ migration and electron transfer, which deteriorates cycling stability. To resolve the shortage of core-shell structure, Sun et al. suggested a core −shell structure with concentration gradient

173.

Sun et al. reported a novel high-

capacity and safe cathode material with an average composition of Li[Ni0.68Co0.18Mn0.18]O2, in which each particle consists of bulk material surrounded by a concentration-gradient outer layer, as illustrated in Figure 18. This material shows not only a very high reversible capacity of 209 mA h g-1 based on the particle bulk composition of Li[Ni0.8Co0.1Mn0.1]O2, but also excellent cycling and safety characteristics, which are attributed to the stability of the concentration-gradient outer layer and the surface composition of Li[Ni0.46Co0.23Mn0.31]O2.

Figure 18 (a) Schematic diagram of positive-electrode particle with Ni-rich core surrounded by concentrationgradient outer layer. (b) Cycling performance of half cells based on Li[Ni0.8Co0.1Mn0.1]O2, Li[Ni0.46Co0.23Mn0.31]O2 and concentration-gradient material cycled between 3.0 and 4.4 V at 55 C by applying a constant current of 0.5 C rate (95 mA g-1). (c) SEM photograph and (d) electron-probe X-ray micro-analysis (EPMA) line scan of the final lithiated oxide Li[Ni0.64Co0.18Mn0.18]O2. Reprinted with permission from ref 173. Copyright 2009 Springer Nature.

Structural mismatch could be mitigated by nano-engineering of the core-shell material, where the shell exhibits a concentration gradient. However, because of the short shell thickness, the manganese concentration at the outer layer of the particle is low; therefore, its effectiveness in stabilizing the surface of the material is weak, especially under high voltage charging. Sun et al.



proposed a full concentration gradient (FCG) structure

168.

The concentration of nickel decreases

gradually from the center towards the outer layer of the particle, whereas the concentration of manganese increases gradually so that the manganese-rich and nickel-poor outer layer can stabilize the material, especially during high-voltage cycling. Besides, the concentration of transition metal is continuously varied from the bulk to the surface, which would minimize the lattice parameter mismatch at the interface between the core and the shell. This material can deliver a specific capacity of up to 215 mA h g-1 at a 0.2 C-rate in the voltage range of 2.74.5 V with outstanding cycling stability in a full-cell configuration. In addition, it also demonstrated excellent capacity retention over 1000 cycles at both 25 °C and 55 °C. This material based on the full gradient approach can lead to the rational design and development of a wide range of functional cathodes with better rate capability, higher energy density and better safety characteristics (Figure 19).

Figure 19. (a) Schematic diagram of the full concentration gradient (FCG) lithium transition-metal oxide particle with the nickel concentration decreasing from the center towards the outer layer and the concentration of manganese increasing accordingly. (b) SEM mapping photograph of Ni, Co and Mn within a single particle for the lithiated material. (c) Cycling performance of half-cells using the FCG, inner composition (IC, LiNi0.86Co0.10Mn0.04O2) and outer composition (OC, LiNi0.70Co0.10Mn0.20O2) materials cycled between 2.7 and 4.5 V versus Li+/Li using a constant current of C/5 (about 44 mA g-1). (d) Electron probe micro-analysis (EPMA) line scan of the integrated


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atomic ratio of transition metals as a function of the distance from the particle center to the surface for the lithiated material. Reprinted with permission from ref 168. Copyright 2012 Springer Nature.

3.5 Electrolyte Electrolyte is an indispensable component in rechargeable batteries, which serves as a carrier ionconductive and electron insulating medium facilitating ion transport between cathode and anode. In LIBs, an electrolyte must have high Li-ion conductivity, high thermal stability and low flammability, and form a stable SEI on the anode. Electrolyte is also very important for the cathode materials in oxidiation stability and surface reactions, especially at high voltage. The electrochemical windows of the electrolytes represent the energy separation of the LUMO and HOMO. A high-voltage stable electrolyte means all the components of electrolytes (solvents, lithium salts, and additives) must simultaneously have lower energy level of HOMO than cathode potential (μc). However, during lithium deintercalation process, μc would shift down to the lower state than the HOMO of electrolyte, resulting in the electron transfer from the electrolytes to the cathodes, and lead to the oxidation of the electrolytes. Therefore, the thermodynamically stable cathode/electrolyte interface requires the LUMO energy level of electrolyte always lower than μc during deintercalation process. Goodenough suggests that a surface film exists on the cathode/electrolyte interface resulting from the reaction between cathode and electrolyte, which is called cathode electrolyte interface (CEI) film and results in kinetic-limited interface stability174. There is an increasing amount of evidence indicating that this protective layer, which possesses similar physicochemical properties as SEI on the surface of anode materials, does indeed exist on the surface of cathode material and also has a significant influence on the electrochemical performance of the battery materials

175, 176.

This reduces the

further side-reactions between cathode material and electrolyte and maintains the stability of the cathode materials at high working voltage, resulting in an excellent performance. Thus, both the lowered HOMO of electrolyte and the stable CEI are crucial for achieving high-voltage TCMs/electrolyte compatibility. 3.5.1 Solvent Electrolyte components provide the prerequisites of interfacial structure and chemistry. So, exploiting new electrolyte components is imperative to obtain high TCMs/electrolyte compatibility at high voltage. New solvent have been developed for the high-voltage electrolyte, including fluorinated carbonates 177-179, sulfones 180-182, nitriles 183, 184, etc. However, these solvents often have



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problems such as poor reductive stability, high viscosity, and lack of stable SEI formation on anode. As a consequence, other co-solvents and/or additives need to be added to, such as cyclic carbonates. Fluorinated carbonates have lower HOMO and LUMO energies than their nonfluorinated counterparts due to the strong electronegativity of fluorine

185.

Therefore, fluorinated carbonates

have higher oxidation potentials and higher reduction potentials. They can be used in high-voltage battery or used as SEI film-forming additives for anodes. Xia et al. reported a fluorinated electrolyte mixture, containing 1 M LiPF6/fluoroethylene carbonate:bis (2,2,2-trifluoroethyl) carbonate (1:1 w:w) with different additives, which exhibited promising cycling and storage performance in Li(Ni0.4Mn0.4Co0.2)O2/graphite pouch type Li-ion cycled between 2.8 V and 4.4 V (Figure 20ad) 177.

They also compared the sulfolane-based and ethylene carbonate-based electrolytes for high-

voltage Li-ion cells containing state-of-the-art electrolyte additives, as shown in Figure 20e, f 180. It showed that the combination of vinylene carbonate and triallyl phosphate as electrolyte additives in sulfolane:ethylmethyl carbonate electrolyte yielded cells capable of better performance during tests to 4.5 V than cells with ethylene carbonate-based electrolytes. These results suggest that sulfolanebased electrolytes may be promising for high-voltage Li-ion cells.

Figure 20. NMC442/graphite pouch cells cycled with clamps at C/2.4 (100 mA) between 2.8 and 4.5 V at 40 ± 0.1 C with different concentrations of prop-1-ene-1,3-sultone (PES) in fluoroethylene carbonate (FEC): bis(2,2,2trifluoroethyl) carbonate (TFEC) electrolyte. (a) Capacity and (b) DV, both plotted vs. cycle number. Panels (c) and (d) compare the PES:FEC:TFEC cell results to those for cells with EC:EMC-based electrolytes as indicated in the legend. Reprinted with permission from ref 177. Copyright 2016 Elsevier. (e) The cycle number when the cell capacity


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reaches 180 mA h (80% retention) for NMC442/graphite pouch cells in sulfolane (SL): ethylmethyl carbonate (EMC) electrolyte with different additive sets as indicated. Comparative data for cells with EC:EMC based electrolytes are also shown. The cycle number has been arranged from “best” at the left to “worst” at the right. (f) A “radar” or “spider” plot which compares the best additive blends. Reprinted with permission from ref

180.

Copyright 2016

Elsevier.

Nitrile-based electrolytes have been reported as additives or co-solvents for high-energy LIBs due to their thermal stability and compatibility with high-voltage cathodes. Farhat et al. reported an electrolyte with adiponitrile (AND) and lithium bis(trimethylsulfonyl)imide (LiTFSI)

186.

It has a

high ionic conductivity over 2 mS cm-1 at room temperature. The Li4Ti5O12/NCM battery cycled between 1.6 V and 2.8 V using this electrolyte exhibited good performance. It has a specific cell capacity of 170 mA h g-1 at low rates (C/20) and even at high discharge rates (2 C), the battery recovers more than 75% of its initial capacity. The excellent stability of the electrolyte permits at least 200 cycles to be performed with a coulombic efficiency close to 100% at each cycle. The oxidative decomposition of nitrile leads to the formation of a cathode interphase, which is responsible for the high-voltage stability of nitrile-containing electrolytes. Despite these remarkable advantages, nitrile-based electrolytes have not been extensive applied in batteries. This is primarily due to its poor reductive stability. Nitrile spontaneously reacts with lithium metal, so graphite and lithium metal electrode does not work reversibly in nitrile-based electrolytes. This incompatibility can be resolved by using high concentration electrolyte or adding SEI-forming additives to produce protect layer on anode 183, 187. 3.5.2 Electrolyte additives Electrolyte additives using in the high-voltage batteries are oxidized predominantly on the cathode electrode surface, similar to how anode additives function, and then form a stable interface layer on the cathode surface that suppresses the side reactions by preventing direct contact between electrolyte and electrode. This is more facile and economic than the surface coating method because the interface layer is formed in situ by adding a small amount of additives into electrolyte. Some common SEI-forming additives on graphite and silicon, such as vinylene carbonate (VC)

188,

fluoroethylene carbonate (FEC) 189, vinyl ethylene carbonate (VEC) 190, and LiBOB 191, 192, LiBF4 193,

show excellent compatibility with layered transition-metal oxide cathodes as well. Besides,

additives consisting of sulfur (e.g., prop-1-ene-1,3-sulton (PES) and methylene methane disulfonate



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(MMDS), divinyl sulfone (DVS)) 188, 194-196, phosphorous (e.g., tris(trimethylsilyl)phosphite (TMSP or TTSPi)) 188, 194, 195, 197, nitrogen (e.g., succinonitrile (SN))197 and boron (e.g., trimethoxyboroxine (TMOBX), tris(trimethylsilyl)borate (TMSB))

191, 198-200

functional groups used in combination

have been reported. Figure 21 compares a large diversity of additives and their combinations in LiNi1/3Co1/3Mn1/3O2 pouch cells in terms of impedance increase, charge end-point capacity ‘slippage’, Coulombic efficiency, and gas production (represented by ‘figure of merits’)

194.

It is

realized that the combination of VC and/or prop-1-ene-1,3 sultone (PES), a sulfur containing additive, like methylenemethane disulfonate (MMDS), together with either tris(-trimethlysilyl)phosphate (TTSP) and/or tris(-trimethyl-silyl)-phosphite (TTSPi) as additives in the electrolyte provide excellent performance of the LiNi1/3Co1/3Mn1/3O2 pouch cells. However, the reason that these blended electrolyte additives work well is still unknown. It is hoped that those highly skilled in surface science and in theory will help unravel these mysteries. In this literature, ultra high precision charger (UHPC), storage, EIS and gas evolution measurements provide a more rapid and better way of screening additives than the simple and common used charge-discharge cycling experiments.

Figure 21. Electrolyte additives for high-performing Ni-rich layered oxides: a systematic survey of several dozen additives for LiNi1/3Co1/3Mn1/3O2/graphite pouch cells in terms of impedance increase, charge end-point capacity


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‘slippage’, Coulombic inefficiency, and gas production. Reprinted with permission from ref 194. Copyright 2014 The Electrochemical Society.

Ma et al. reported a systematic study of some promising electrolyte additives in Li[Ni1/3Mn1/3Co1/3]O2/graphite, Li[Ni0.5Mn0.3Co0.2]/graphite and Li[Ni0.6Mn0.2Co0.2]/graphite pouch cells

188.

The different properties such as charge end-point capacity slippage, capacity decay,

coulombic efficiency, impedance change within cycling, gas evolution as well as voltage drop during “cycle-store” testing were also evaluated and compared as illustrated in Figure 22. Although there is no absolute “winner” among these additives or combinations, it is clear that “PES211”(2% PES + 1% methylene methanedisulfonate (MMDS) + 1% tris(trimethylsilyl) phosphite (TTSPi)) and 0.5% pyrazine di-boron trifluoride (PRZ) + 1% MMDS yield significant advantages in most aspects of cell performance compared to the other electrolyte additives. Unfortunately, the effect mechanisms of these additives and blend additives are unknown, which need more efforts to elucidate.

Figure 22. Radar plots summarizing the effects of selected electrolyte additives (combinations) on NMC111/graphite (a), NMC532/graphite (b) and NMC622/graphite (c) pouch cells studied using UHPC and the “cycle-store” procedure to 4.4 V. The axes are normalized to the worst value being equal to 100% and they consist of the average coulombic inefficiency (CIE) (from 11 to 15 cycles), the average charge end-point capacity slippage (from 11 to 15 cycles), the impedance (Rct) after ultra high precision charger (UHPC) cycling, the gas evolution



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during UHPC cycling, the capacity loss after 35 “cycle-store” cycles, the impedance after the whole “cycle-store” process, the voltage drop at 35 “cycle-store” cycle and the gas evolution during the whole “cycle-store” process. Values closest to the center of the radar plot are best. Reprinted with permission from ref 188. Copyright 2015 Elsevier.

3.5.3 Ionic liquids Ionic liquids (ILs), defined as room-temperature molten salts with melting point below 100 C, are composed mainly of organic cations and (in)organic anions ions. They offer a unique series of physical and chemical properties, such as nonvolatility, high thermal stability, and wide electrochemical window 201-204. Nonetheless, only a few ILs have been used as solvents or additives for electrolytes in LIBs with TCMs, because most of the ILs have disadvantages, such as poor electrochemical reductive stability, poor wetting, relatively high viscosity, and co-intercalation into the graphitized carbon-based anode

205-207.

Lee et al. reported two types of monocationic ILs 1-

butyl-3-methylpyrrolidinium hexafluorophosphate (Pyr IL) and 1-ethyl-3-methylimidazolium hexafluorophosphate (IMI IL) as an additive in 1.15 M LiPF6 (EC/EMC=3/7 v/v) electrolyte solution

205.

The electrochemical measurements showed that the Pyr IL additive has better

compatibility with NCA cathode active material than IMI IL additive. LiNi0.80Co0.15Al0.05O2 /carbon full-cell with Pyr IL additives at 1 wt% and 3 wt% deliver better cycleability than others, with the retention ratios of 93.62% and 92.8%, respectively. At elevated temperature (55 C), only 1 wt% Pyr IL additive is giving stable capacity retention ratio of 80.74%. 3.5.4 Polymer electrolyte Though many strategies have been developed based on liquid electrolyte to address the side reaction between electrolyte and TCMs, the safety issues still exist when the batteries operate under high voltage, attributing to the oxygen release. Therefore, the polymer electrolyte with limited liquid or without liquid is considered to be an alternative to balance the electrochemical performance and safety of the batteries with high operating voltage TCMs

208-215.

Cui et al. reported a composite

membrane consisting of poly(ethylene terephthalate) (PET) nonwoven fabric incorporated with poly(ethyl α-cyanoacrylate) as polymer matrix of gel polymer electrolyte (GPE)

215.

The ionic

conductivity of the electrolyte based on this composite membrane saturated with a liquid electrolyte (LE) of 1M LiPF6 in carbonate solvents is 2.54 mS cm−1 at room temperature, and the stable electrochemical voltage window can be up to 4.7 V. Its electrochemical performance is evaluated by using full cell with Li[Ni0.5Co0.2Mn0.3]O2 as cathode and graphite as anode. The GPE can


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suppress the side reaction between NMC cathode and electrolyte and restrain the metal dissolution at high operating voltage as shown in Figure 23. Compared with liquid electrolyte, the capacity retention of cell using GPEs increased from 20 to 77% at potential range of 3.0–4.6 V. Though polymer electrolyte exhibit higher mechanical and thermal stability than LEs, there are limited kinds of polymer electrolytes applied in high-voltage TCMs based LIBs for their lower oxidation voltage. It needs more effort to explore new polymer electrolyte with wider electrochemical window. Based on the above mentioned thermodynamic and kinetic-limited interface stability, there are three main strategies to improve the antioxidative capability of polymer electrolyte: (1) molecular designing of polymer, such as copolymerization or grafting with chemical inert main chain structure and the polar group with low HOMO energy level 216; (2) utilizing the intermolecular interaction between polymer and lithium salt to lower HOMO of electrolyte

217-219,

such as ion-dipole and Lewis acid-based

interaction; (3) using lithium salts or additives, which are oxidized predominantly on the cathode surface to form stable interface layer and protect the polymer electrolyte from further oxidization 220-222.

Figure 23. Cycle performance of NMC/G full battery in the cut-off range of 3.0–4.3 V (a) and 3.0–4.6 V (b); charge– discharge curves of batteries cycling in the cut-off range of 3.0–4.6 V in GPE and liquid electrolyte (c); Cycle



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performance of NMC/G in GPE soaked in LiBOB electrolyte at the cut-off range of 3.0–4.6 V (d). Reprinted with permission from ref 215. Copyright 2015 Elsevier.

3.5.5 Solid state inorganic electrolytes Besides the gel polymer electrolytes with somewhat liquid, solid state electrolytes containing no liquid with high safety and wide electrochemical window represent another promise alternative to liquid electrolyte using in the high-voltage cathode battery system

223.

The solid state electrolyte

would provide high safety and high energy density of lithium metal batteries with high operating voltage TCMs. Solid state inorganic electrolytes include oxides, such as perovskite Li3yLa2/3−yTiO3 (LLTO) 224, garnet Li7La3Zr2O12 (LLZO) 225, LiSICON-type phosphates 224, as well as sulfide solid electrolytes such as Li10GeP2S12 (LGPS) and thio-LiSICON 226, 227 have been widely researched for many years. Among all the inorganic solid electrolytes, the sulfide solid electrolytes show the highest conductivity at room temperature (10-3–10-2 S cm-1)

226, 227.

Hirokazu et al. reported

LiNi1/3Co1/3Mn1/3O2 applied to all-solid-state batteries using the 80Li2S19P2S51P2O5 (mol %) solid electrolyte. The experimental cell exhibited a first discharge capacity of 115 mA h g-1 at a current density of 0.064 mA/cm2 and retained a reversible capacity of 110 mA h g-1 after ten cycles during 2.5–4.4 V 228. The sulfide electrolyte appears to be incompatible with LiNi1/3Co1/3Mn1/3O2 because the decomposition reactions and space charge layer (SCL) at the interface. Coating LiNi1/3Co1/3Mn1/3O2 particles with Li4Ti5O12 was effective for decreasing the interfacial resistance and improving the electrochemical properties such as rate performance. However, the interface issues of high-voltage layer cathode materials and solid electrolyte remain serious. As we mentioned before, the heterogeneous change of lattice parameters during charging/discharging would cause the microcrack, and meanwhile lead to poor contact of cathode materials and solid state electrolyte, which results in a large interface resistance. Besides, the SCL caused by interaction of cathode and electrolyte provides high resistance and blocks the transportation of lithium ions. Especially under large current density, the batteries suffer from large polarization and rapid capacity degradation. Meanwhile, the composite cathode with solid electrolyte will decrease the energy density of the batteries. The use of solid state electrolyte just suppresses the corrosion of liquid electrolyte and dissolution of metal element but cannot improve the phase transformation of TCMs. So the solid electrolyte cannot give full play to the advantages of high-voltage TCMs and it still need more efforts to solve these problems.


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4. REMAINING CHALLENGES Although a variety of the cathode modification strategies considerably enhance the structure and interface stability and achieve higher energy density and cycle performance, the application of highvoltage TCMs in EVs is still severely challenged by air sensitivity, safety issues and lower rate capability. The residual alkali on the surface of TCMs makes it quite easy to absorb water and then increases difficult in the preparation of slurry. More importantly, the residual alkali compounds tend to decompose at high voltage and produce gas during battery cycling. The exothermic reactions and oxygen release at high voltage will also increase the risk of thermal runaway. 4.1 Air storage properties The NMC and NCA materials, especially with high Ni content, are sensitive to the CO2 and H2O when exposed in the air atmosphere 229, 230. Active oxygen species that are generated at the particle surface due to the reduction of Ni3+ to Ni2+ are prone to react with trace amounts of CO2 and H2O in air, forming a layer of residual lithium compounds on the surface that mainly consist of Li2CO3 and LiOH. Besides, a NiO layer in contact with the bulk material also forms due to the depletion of Li and loss of lattice oxygen

231, 232.

Both layers delay Li+ diffusion and cause large irreversible

capacity during initial cycle. The fresh LiNi0.70Co0.15Mn0.15O2 cathode revealed clean surface composed of the primary particles, whereas LiOH and Li2CO3 were observed in the LiNi0.70Co0.15Mn0.15O2 after aging in air for 3 months

233.

The LiOH on the cathode increases the

powder pH value, causing the gelation of the slurry during the electrode fabrication process. The Li2CO3 significantly promotes the gas evolution and increase the moisture of the cathode powder. The residual lithium compounds were electrochemically oxidized at ~ 4.1 V, generating the CO2 gas. Furthermore, at a high charged state, the residual lithium compounds give rise to the CO2 gas evolution by reacting with electrolyte 234, 235. The initial gas evolution is proportional to the charge cut-off voltage because of enhanced nucleophilic reaction of the electrolyte and the formation of large amount of HF at higher voltage. Recently, there is an new viewpoint that the overwhelming majority of CO and CO2 evolution comes from oxidation of residual lithium compounds and does not originate from electrolyte decomposition, up to 4.8 V vs Li/Li+

236.

During storage, the

spontaneous reaction of Ni3+ to Ni2+ is likely to take place, causing a loss in structural ordering. This surface deterioration leads to a large increase in resistance, impeding the charge transfer and the mobility of lithium ions and electrons for the electrochemical redox reaction, which would further



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result in the fast decay of electrochemical properties after storage in air for a period of time. It suggests that the residual lithium compounds should be removed from the surface of TCMs particles to prevent gas generation and performance deterioration

232.

There are three main strategies to

handle this problem. The first is to isolate the TCMs from H2O and CO2, such as storing TCMs in drying room. It needs an extra cost to build drying room. The second is to recover the initial condition by a post-treatment strategy, such as washing by water. However, the washing process makes TCMs more chemically sensitive than non-washed materials and give rise to severe capacity fading at a high temperature of ≧45C 237. The last strategy is to introduce surface coating 238, 239. The coating precursor could react with the residual lithium compounds during annealing process, thereby diminishing these compounds and producing coating layer. These methods have improve the cycling performance of TCMs but increase the cost and process time, so it needs to develop an simple and economical method to improve storage properties of TCMs. 4.2 Safety issues LIBs in a high charged state can be very dangerous if exothermic reactions start between the active materials and organic electrolytes. The triggering mechanism is related to the cathode structural instability, which results in O2 and CO2 release. The oxygen species released from the cathode framework are so reactive with the electrolytes, which contribute greatly to the thermal runaway. Detecting the onset temperature of phase transformation and O2/CO2 release and the amount of gas release are the common methods to evaluate thermal stability of charged cathode materials. Many techniques have been applied to study the thermal stability of charged cathode materials, such as time-resolved and high-temperature X ray diffraction (XRD) thermogravimetric analysis (TGA)/differential scanning calorimetry (DSC) spectroscopy (MS) and accelerating rate calorimetry (ARC)

244-246.

242,

240, 241,

243,

mass

Yang et al. combined time-

resolved synchrotron XRD and MS to explore the structural changes and the evolution of gas that occurs during the thermal decomposition of charged LixNi0.8Co0.15Al0.05O2 247. It indicates that the evolution of both O2 and CO2 gases are strongly related to phase transitions occurring during thermal decomposition, specifically from the layered structure to the disordered spinel structure, and finally to the rocksalt structure. The SOC significantly affects both the structural changes and the evolution of oxygen as the temperature increases: the more extensive the charge, the lower the temperature of the phase transitions and the larger the oxygen release. They also compared the structural changes


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of overcharged LixNi0.8Co0.15Al0.05O2 and LixNi1/3Co1/3Mn1/3O2 during heating by TEM and X-ray techniques 248. The scheme of Figure 24 illustrates that the NCA particles consist of a rhombohedral core, a spinel shell, and a rock-salt structure at the surface at room temperature. During heating, the

surface structures propagate towards to the core, and the rocksalt phase becomes the dominant structure at high temperature. NCM particles consist of O1 structure at the surface at room temperature. During heating, this phase transforms into the Co3O4 -type-spinel structure and stays as such to high temperature. The right panel illustrates migration of Co cations to the tetrahedral sites and the unchanged Mn at original octahedral sites. These results suggest that Mn plays a major role in suppressing the formation of the rock-salt structure on the surface, and the combined effects of Mn and Co help to form and maintain the Co3O4 -type spinel and suppress the transformation to the rock-salt structure during heating. Therefore, increasing the Mn content on the surface of the particles could effectively improve the thermal stability of the layered cathodes, while retaining the high-energy density by keeping high Ni concentration in the bulk. The thermal stability of the active material is also heavily dependent on the interface between active materials and electrolytes. Surface coating is an effective way to improve the thermal stability, which can inhibit not only the formation of by-products from electrolyte decomposition reactions but also the spread of cation disordered-phases. However, the structure instability and oxygen release is still an important problem of the TCMs when they charged at high voltage. Regardless of the intensive efforts made to understand thermally stimulated structural transitions TCMs, the basic correlation has remained unclear between the surface/interfacial structure and chemistry. The safety issue promises to be solved by the replacement of flammable liquid electrolyte by thermostable solid-state electrolytes. However, much thermal analysis experiments are needed to evaluate the thermal stability of TCM-based solid state batteries.



Figure 24. Schematic of the mechanism of thermal decomposition of the overcharged (a) LixNi0.8Co0.15Al0.05O2 and (b) LixNi1/3Co1/3Mn1/3O2 cathode during heating. The left side shows phase propagation from the surface to core for the overcharged particle. The right side shows the changes in crystal structure and cation distribution. Reprinted with permission from ref 248. Copyright 2013 Wiley.

4.3 Rate capability Rate capability is an important criterion to evaluate cathode materials and is affected by several factors, such as ratio of metal elements, cation mixing, and surface characters. In TCMs, Co ions are beneficial for both electrical and ionic conductivity of the materials, leading to good rate capability. The presence of electrochemically inactive Mn in the oxide lattice ensures structural and thermal stability but promote more Ni2+ formation (charge neutrality), which deteriorates the rate capability. Al ions function in a similar way to Mn ions. So the rate capability of TCMs is dependent on the ratio of these elements. Compared to the well-ordered layer structure, the disordered phase has a higher activation energy barrier for lithium diffusion owing to its smaller distance between the slabs, and also a lower Li+ diffusivity caused by the transition metal in the lithium layer. Therefore, the rate capability of the TCMs decreases with the increasing cation mixing. Residual lithium compounds on the materials surface have low electronic and ionic conductivity, impeding the mobility of lithium ions and electrons for the electrochemical redox reaction, resulting in low rate capability. The surface phase transformation caused by release of O2 and corrosion of electrolyte also impacts the rate capability of materials. Compared to the well-ordered phase, the spinel and/or rock-salt structures possess higher activation energy barrier for Li+ diffusion, resulting in sluggish Li+ kinetics during redox reactions, thus leading to poor rate performance of TCMs. These situations are more serious at high charged voltage as we discussed before, so the rate


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capability would deteriorate as the voltage increases. There are several strategies to increase rate capability of TCMs. The most prevalent method is surface coating, which can either act as electronic/ionic conductivity facilitator or shield the materials from air atmosphere and organic electrolyte, alleviating the surface phase transformation 133, 249.

Surface coating increases rate capability by decreasing the charge-transfer resistance,

increasing particle conductivity, and/or altering the thickness and composition of surface layer. Element doping is another simple approach to enhance rate capability 71, 250. Certain dopants with larger ionic radius can increase the Li layer spacing, which reduces the energy barrier and facilitates Li+ diffusion. Besides, element doping is benefit for the structural stability and thus in favor of high Li+ diffusivity. There are some additional methods to enhance the rate capability, such as producing suitable size particles to obtain shorter Li+ diffusion paths materials to form a composite

252.

251;

and blending TCMs with other

But, recent TCMs still cannot meet the consumer demands for

fast charge, so it needs more effort to realize high rate capability. 5. CONCLUSION AND PROSPECT In this review, we discuss the recent progress in high-voltage TCMs and the related key problems and solutions for high energy density LIBs used in PHEV/EVs. The issue of high-voltage electrochemical operation of TCMs materials lies mainly in the stability of the electrodes and electrolytes as well as their interfaces. Structure variations and exfoliations, electrolyte parasitic reaction, oxygen loss, and microcrack are major degradation processes and ultimately lead to irreversible collapse of TCMs structures, loss of electronic contact, increase in internal resistance, and reversible capacity degradation. Some strategies have been proposed, including elemental doping into the crystal framework, coating on the surface, special structure/composition design, utilizing new electrolytes, electrolyte additives, etc. Element doping is a simple method to adjust the lattice parameter at atom scale, stabilizing intrinsic structure at high operating voltage. Most of the doping elements are electrochemical inactive which would decrease the capacity of TCMs, so the doping element content should be considered to balance the capacity and cycling performance. Surface coating is an effective method to suppress side reactions between TCMs particles and electrolyte, stabilize the interface structure and suppress structure transformation and the dissolution of transition-metal ions at high working voltage. However, apart from electronic and ionic conductive coating materials, the coating layers always decrease the capacity of TCMs more or less.



Moreover, the thickness should be reasonable and suitable, otherwise it cannot provide a sufficient protection or lower the capacity and rate performance. The approach to control the thickness of coating materials by thin film technology can increase the cost of TCMs production. Though the co-modification of surface coating and element doping combines their advantages, their disadvantages still exist. Many additives were reported in high-voltage electrolyte and improve the cycling performance, but the effect mechanism still a mystery. More attention should be focused on surface/interfacial stability by studying the SEI layer-forming mechanisms in future work. The surface and interface chemistry properties play a critical role in determining the structural and electrochemical stability of TCMs in high-voltage LIBs. Electrochemically stable electrolytes and/or electrolyte additives, as well as conductive additives, binders, and current collectors stable at high voltage need to be identified to enable such high-voltage cathode materials. Range anxiety and safety are two major challenges for the long-term development of long-range EVs powered by liquid organic electrolyte-based LIBs. Solid state batteries are believed to overcome the above challenges and have been well laid out in the roadmap of power battery development by several countries. In the nearly future, high-voltage TCMs-based solid state batteries will be a hot but challenging research topic in the field of long-range EVs. To achieve the commercial application of high-voltage TCMs, our future work will not only focus on the TCMs themselves but also exploring solid state electrolyte with wide electrochemical window and high ionic conductivity and optimizing electrode/solid electrolyte interfacial compatibility. The solid sulfide electrolyte has high ionic conductivity and exhibits plastic deformation under mechanical pressure which makes it an intimate contact with cathode material. However, the SCL caused by the Li+ chemical potential difference between oxide cathode and solid sulfide electrolyte can impede Li+ transport at the cathode/electrolyte interface, resulting high polarization and capacity degradation. The rigid ceramic nature of solid oxide electrolyte makes poor contact between cathode and solid electrolyte, and the interface will deteriorate when the lattice parameters and volume significantly change at high charge state. So, how to optimize the TCMs/ inorganic solid electrolyte interface is the foremost task. For the solid polymer electrolyte, its interface compatibility with cathode is superior to the inorganic solid electrolyte. Moreover, due to the poor processability and high cost of inorganic electrolytes, we think polymer electrolytes will be more attractive in the commercialization of solid state batteries. But the poor antioxidative capability is the main


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disadvantage. There are two factors that can make polymer electrolytes provide good stability under high voltage operation. One factor is the wide electrochemical stable window, which means polymer electrolyte having lower energy level of HOMO than cathode potential and resulting in thermodynamically stable cathode/electrolyte interface. The other factor is the kinetic-limited stability achieved by forming the chemical passivation layer at the cathode/electrolyte interface when the electrode electrochemical potentials μC is outside the electrochemical window of polymer electrolyte. In this field, our group has explored a series of polymer electrolytes with wide electrochemical window by designing the chemical structure of polymers and tuning the intermolecular interaction between polymer and lithium salt/additives, which is summarized in Figure 25

253-261.

These polymer electrolytes would have promise practical use in high-voltage

TCMs.

Figure 25. Summary of the high performance solid/gel polymer electrolyte reported by our group. Reprinted with



permission from ref

255.

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Copyright 2017 American Chemical Society. Reprinted with permission from ref

Copyright 2016 Royal Society of Chemistry. Reprinted with permission from ref

258.

256.

Copyright 2017 American

Chemical Society. Reprinted with permission from ref 259. Copyright 2017 Royal Society of Chemistry. Reprinted with permission from ref 260. Copy Copyright 2017 Royal Society of Chemistry. Reprinted with permission from ref 261.

Copyright 2018 Royal Society of Chemistry.

In summary, the high-voltage TCMs have been understood deeply these years from the electrochemical properties, structure transformation, interface reactions, failure mechanism and modification approaches. Though many strategies have been proposed to improve the high-voltage TCMs in lithium batteries, they still have many challenges related to structural instability, storage degradation, surface side reactions, safety issues, and microcrack and there is a long and winding road ahead for the high-voltage TCMs to reach the commercial applications. Even so, it is believed that through continuous scientific efforts the high-voltage TCMs would be applied in the EVs and relieve the range anxiety of EVs in the near future. ASSOCIATED CONTENT AUTHOR INFORMATION Corresponding Author *E-mail: [email protected] and [email protected] Notes The authors declare no competing financial interest. ACKNOWLEDGMENT This original research was supported by funding from National Science Fund for Distinguished Young Scholars (Grant No. 51625204), the National Key R&D Program of China (No. 2018YFB0104300), National Natural Science Foundation of China (No. U1706229). REFERENCES (1) Armand, M.; Tarascon, J. M. Building better batteries. Nature 2008, 451, 652-657. (2) Ren, G.; Ma, G.; Cong, N. Review of electrical energy storage system for vehicular applications. Renewable Sustainable Energy Rev. 2015, 41, 225-236. (3) Yang, Z.; Zhang, J.; Kintner-Meyer, M. C.; Lu, X.; Choi, D.; Lemmon, J. P.; Liu, J. Electrochemical energy storage for green grid. Chem. Rev. 2011, 111, 3577-613. (4) Scrosati, B.; Hassoun, J.; Sun, Y.-K. Lithium-ion batteries. A look into the future. Energy Environ. Sci. 2011, 4, 3287-3295. (5) Thackeray, M. M.; Wolverton, C.; Isaacs, E. D. Electrical energy storage for transportation-


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Identifying and Addressing Critical Challenges of High-Voltage

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