Degradation Mechanisms and Mitigation Strategies of Nickel-Rich NMC-Based Lithium-Ion Batteries

29 Oct.,2022


700mm UHP Graphite Electrode

Ni-rich NMC-based materials are some of the most promising cathode candidates for next-generation LIBs due to high capacities and large voltage windows. Despite impressive progress in the development of advanced materials and fabrication processes, the development of high-capacity electrodes with long-term stability and prolonged cycle lifespans remains an important challenge. To address this, the understanding of underlying degradation mechanisms presented in this review can serve as a starting point to learn and improve the performance of Ni-rich NMC-based cathodes. In addition, the information included in this review can provide valuable insights into the creation of new devices with improved efficiency and productivity. And for the sake of completeness in the understanding of degradation mechanisms in LIB full cells, parasitic reactions on NMC-based cathode surfaces, the aging of graphite anodes as well as SEI formation and aging are also briefly presented in this section.

Ni-Rich NMC-Based Cathodes

The degradation of Ni-rich NMC LIBs involves different types of capacity loss, including initial capacity loss (ICL), sudden capacity loss (SCL) and gradual capacity loss (GCL), all of which are accompanied by impedance growth. And in the case of GCL, which is the most frequently investigated process, a major cause is the loss of Li and transition metals on cathode surfaces. In general, the majority of degradation mechanisms for Ni-rich NMC-based cathodes can be indexed to the same mechanism independent of Ni/Mn/Co ratios because they share similar crystal and microspherical structures. Here, adjustments (typically increases in Ni) to the relative amounts of Ni, Mn and Co can result in augmented performances as well as various negative features. For example, enhanced mass-specific capacity can be achieved but at the expense of rate capability and structural stability. In addition, researchers have also reported higher electronic conductivity and reduced polarization [40] in which the varied polarization as compared with the portion of Ni is caused by the higher eg orbital of Ni over the t2g orbital of Co [41]. Researchers have also shown that the reduced Mn content can result in faster degradation through surface reconstruction [42] and that the increased Ni content can result in increased cation mixing due to the bulk diffusion of Ni2+ into the Li-layer and increased parasitic reactions as the valence of surface Ni increases toward highly reactive Ni4+. Overall, cathode performances can be impacted by various factors that interact with each other, and therefore, this section will discuss degradation mechanisms in detail.

Surface Degradation During Cell Operation

Lattice Expansion/Contraction

NMC oxide possesses an \( R\bar{3}m \) structure (rhombohedral symmetry) with a Li-layer on the 3a site, an NMC layer on the 3b site and an oxygen layer on the 6c site [7] and can usually be indicated in the splitting of (110) and (108) peaks and of (006) and (102) peaks. For example, the expansion and contraction of the c-axis (14.21 Å) of NMC811 lattices can be observed by using in situ XRD measurements in the cell voltage windows of 3.0–4.0 V and 4.0–4.4 V, in which (003) peaks can decrease from 18.96° (3.0 V) and subsequently rise to 19.09° (4.4 V) [7, 10, 43]. This can also be observed for the movement of (104), (015) and (108) peaks. Here, researchers suggest that the expansion of the c-axis is due to the repulsive force generated from MO6 slabs, which are positively charged in a highly delithiated state, whereas the loss of electrons can shrink the radius of transition metal ions to reduce the a-axis [44]. In addition, the (110) peak can shift to higher angles in experiments, suggesting that the a-axis can contract during charging due to the smaller radius of transition metal ions at elevated valences [45]. Researchers have also reported that the a-axis of NMC811 lattices remains constant at potentials exceeding 4.3 V [10] and that a steep drop of the c-axis occurs at potentials exceeding 4.2 V in which the drop becomes steeper as the Ni content increases [7, 20, 43]. Here, researchers proposed that this drop is due to the fact that the repulsion between O2− layers decreases with more covalent M–O bonding at higher delithiated states [46, 47] (Fig. 1a, b). Furthermore, studies have also confirmed two pathways (Fig. 1c2, c3) involved in the delithiation of Ni-rich NMC-based cathodes [40, 48], including oxygen dumbbell hopping and tetrahedral site hopping, in which at the beginning of charging, Li can exit through oxygen dumbbell hopping until a certain degree of delithiation is reached and continue through tetrahedral site hopping due to increased energy barriers as caused by Li–O bonds during delithiation.

Fig. 1

The c-axis (a1), the a-axis (a2) and potential (a3) of an NMC 811 cell as a function of specific capacity. The c-axis (b1) and the a-axis (b2) as a function of cell potential during the second cycle and dQ/dV of the second cycle as a function of cell potential (b3) (Reprinted with permission from Ref. [7]. Copyright © The Author(s) 2015. Published by ECS). c1 Lattice of an NMC layered structure. c2 Tetrahedral site pathway and c3 oxygen dumbbell pathway for Li-ion diffusion in an NMC layered structure (Reprinted with permission from Ref. [48]. Copyright © 2015 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim)

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Surface Reconstruction to Fm3m Rocksalt

Surface reconstruction of the particles of NMC-based cathodes can initiate if delithiation exceeds 70% as shown in NMC111 [49]. In addition, increasing the Ni content can lead to increased portions of Ni3+ (usually Ni ions with higher valences) in the total amount of Ni. Furthermore, elemental Ni partially located in the transition metal layer as Ni2+ (0.69 Å) possesses a similar radius to Li+ (0.76 Å) [50], and Li+ and Ni2+ can exchange positions in a delithiated state through Ni2+ diffusion to the octahedral sites of Li+ (energetically favorable) through a neighboring tetrahedral site (Fig. 1c) [42, 51]. Here, the degree of Ni2+ migration can be evaluated by calculating the c/a ratio of the NMC lattice because this migration reduces the c/a ratio. And based on XRD measurements, ratios of I003 to I104 can decrease to less than 1.2 after significant cation mixing and can therefore serve as an indicator of the degree of cation mixing [52]. Moreover, researchers reported that cation disorder can lower spacing between atomic layers in the lattice of the NMC-based cathodes, hinder Li+ movement and reduce the amount of active Ni and Li [53]. In the case of higher Ni percentage cathodes, Mn4+ can partially be replaced by Ni, which increases the valence of Ni and therefore lowers the possibility of Ni atoms migrating to 3a sites [8]. However, because the total amount of Ni also increases, Ni2+ migration still occurs to a larger extent. To minimize Ni2+ migration, researchers reported that mitigation methods can be used to modify cathodes [54]. Furthermore, a disordered spinel structure (Fd3m) due to non-ideal cation mixing [55] can form if LIBs are charged to high voltages (e.g., 4.8 V) [21, 56,57,58,59,60] and consists of a LiM2O4 spinel structure and a M3O4 spinel structure if Co migration into tetrahedral sites occurs [42] (Fig. 2a). Xiong et al. [61] and Eom et al. [62] also reported that spinel phases can be formed on charged cathodes stored at 90 °C for a week. Moreover, researchers found that migrated Mn and Co possess high valences and smaller radii for movement and more vacancies in the Li-layer in a highly delithiated state [63]. Researchers also reported that long-term cycling to 4.2 V (NMC 622) [64] can result in phase changes to an Fm3m structure consisting of Ni2+, Mn2+ and Co2+ with lowered valence states as compared with a spinel structure [63, 65] and ion-insulating [56, 63, 66] (a cubic rocksalt phase) (Fig. 2b). Here, the conversion from a layered structure to a cubic rocksalt phase (usually confirmed by HRTEM/EELS [9]) can release O2 into the cell, which can react with electrolytes to generate CO2 and result in increased electrode interfacial resistances due to kinetic barriers in the Li+ insertion/desertion process [50]. To address this, researchers reported that NMC811 cathodes pretreated with ramping to 4.5 V limited the growth of the rocksalt layer in subsequent cycles and can serve as a pillar structure to stabilize the lattice and isolate the surface of NMC particles [67]. Furthermore, researchers reported that the oxidation states of Mn and Co decrease in the surface layer/region of NMC811 and NMC442 after cycling [9, 63].

Fig. 2

a Schematic of phase transition and possible TM cation migration pathways in charged NMC cathode materials during thermal decomposition (Reprinted with permission from Ref. [42]. Copyright © 2014, American Chemical Society). b Degradation mechanisms of LiNi0.5Co0.2Mn0.3O2 after cycle tests under two upper cutoff voltage conditions (4.5 V and 4.8 V) (Reprinted with permission from Ref. [56]. Copyright © 2013 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim). c Differential capacity plot of three different NMC/graphite cells recorded at 0.1 C-rate (3rd cycle) with marked phase transformations (Reprinted with permission from Ref. [25]. Copyright © The Author(s) 2017. Published by ECS). d Schematic of H1–H2–H3 phase transformation from the perspective of local environments in different views along the c-axis (up) and the a-axis (bottom) (Reprinted with permission from Ref. [69]. Copyright © 2018, Elsevier B.V. All rights reserved.)

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Phase Transformations

Two hexagonal phases of NMC811 can coexist on the surface of particles if freshly prepared in which the evolution of one phase is irreversible and vanishes after initial charging (with lattice constants remaining at the initial value), whereas the second phase is reversible (except for a small difference between charging and discharging above 4.2 V). Here, the fading of the (110) peak at 65.12° can track the conversion of the region in which the two hexagonal phases coexist and are converted to a single phase [7]. In addition, the broadening of the (003) peak during cycling as recorded by using in situ XRD measurements is of great interest in which researchers [25] suggested that this phenomenon was due to the creation of a second phase in the potential range of 3.8–4.0 V (H2, evolved from the monoclinic phase in contrast to the phase H1 following the stoichiometric composition) that possesses less Li. Researchers also reported that this process occurred earlier with higher Ni portions (Fig. 2c) and that at higher voltages, the conversion from H2 to H3 occurs with impedance growth (volume contraction) only for NMC811 (or NMCs with Ni portions larger than 80%) [25]. Researchers have also found that the phase transition from H2 to H3 can lead to (003) peaks shifting to higher degrees [68] and that Li-reordering (Li diffusing into Ni4+-rich sites) can further convert H3 to H3-2 [69] in which the fast shrinkage of the c-axis at potentials above 4.15 V is a clear indication of this phase transition. Moreover, the generated microregion of NiO2 in the intralayer of the NMC811 lattice as a result of the phase transition from H2 to H3 can lead to the irreversible transformation of the NMC811 structure [45] and this phase transition becomes severer as Ni portions further increase to extreme numbers as indicated by rising peaks of the differential capacity curve centered at 4.15 V [68]. Here, the fading of this peak leads to capacity fade and is referred to as the “deleterious effect.” Furthermore, phase changes to a second hexagonal phase for other NMC-based cathodes have yet to be clarified. (Fig. 2d illustrates crystal lattice evolution during delithiation.)

Surface Stabilizer

In general, NMC-based cathodes with more Mn content tend to possess better cycling stability [13, 70], whereas NMC-based cathodes with more Ni content are believed to be incapable of reaching high cutoff voltages due to the lack of Mn4+ as a structure stabilizer [65]. Apart from Ni2+, structural instability also applies to Mn3+, although with the high Ni content, Mn migration is not considered to be significant. Here, due to low octahedral site stability energy, NMC-based cathodes possess a tendency to form spinel-structured LiMn2O4 on surfaces [71] and Mn4+ possesses a similar tendency to migrate to the Li-layer [33, 71] and decrease capacity. Nevertheless, diffused Mn4+ on particle surfaces can stabilize structures during long-term cycling [50].

Impurities and Parasitic Reactions

Surface Impurities

To maintain a layered structure, excessive Li is used for Ni-rich NMC-based cathodes. However, side reactions of Li with water vapor and CO2 in air can lead to LiOH and Li2CO3 in which exposure to air can result in a reconstructed surface layer 3 nm thick (evidenced by STEM) that thickens if cycled [9]. A surface layer with similar impacts can also form as cathodes come into contact with electrolytes [65]. Here, researchers suggest that this reconstructed layer can be tuned through different synthesis methods of the cathode powder [63] and that in the case of NMC111, surface impurities can also grow on top of cathode particles as resistive films over the same time period for uncycled Ni-rich NMC-based LIBs [35] that contain mainly insulating hydroxides and carbonates as revealed through Raman spectroscopy (Fig. 3a). Here, the circumstances in which Li2CO3 (Eq. 1) or NiCO3 is favored remain unclear [35, 72, 73]. Researchers also reported that high humidity can lead to more hydroxide (e.g., LiOH) on NMC622 [72] and that surface impurities can affect cell impedances in subsequent cycles through reactions with electrolytes to form a Li+ diffusion inhibiting layer.

$$ {\text{Li}}\left( {{\text{Ni}},{\text{Mn}},{\text{Co}}} \right){\text{O}}_{2} + \frac{x}{2}{\text{CO}}_{2} + \frac{x}{4}{\text{O}}_{2} \to {\text{Li}}_{1 - x} \left( {{\text{Ni}},{\text{Mn}},{\text{Co}}} \right){\text{O}}_{2} + \frac{x}{2}{\text{Li}}_{2} {\text{CO}}_{3} $$


Fig. 3

a Raman spectra of fresh and stored NMC111 and NMC811 electrodes before cycling, measured in air within ~ 1 h with the samples not treated prior to measurements (Reprinted with permission from Ref. [35]. Copyright © The Author(s) 2018. Published by ECS). b Positions of various redox couples relative to the top of the oxygen 2p band (Reprinted with permission from Ref. [14], Copyright © 2017, American Chemical Society. Further permissions related to this figure should be directed to ACS). c Proposed electrooxidation and chemical oxidation pathways for ethylene carbonate (EC) and their potential dependence (Reprinted with permission from Ref. [25]. Copyright © The Author(s) 2017. Published by ECS). d Schematic of the different issues facing Ni-rich NMC materials: (1) SEI formation and reactivity of Ni4+; (2) Li/Ni displacement and formation of disordered phases; (3) microcracks in secondary particles (Reprinted with permission from Ref. [13]. Copyright © The Author(s) 2016. Published by ECS)

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NMC Dissolution

The dissolution of NMC materials can also result in capacity attenuation [26, 74] because it can decrease Li+ insertion sites and is triggered by the products of electrolyte decomposition such as HF from the reaction of PF5 with H2O (Eqs. 2, 3) [31]. In addition, higher voltages can accelerate the dissolution of NMC111 due to the release of more acidic components as a result of rapid electrolyte decomposition [75]. Despite these findings, effects on Ni-rich NMC have yet to be reported. The dissolution of NMC materials can also cause a concurrent issue involving the production of resistive MFx as a side product on the surface of NMC particles. Moreover, the dissolution of transition metals can pose a threat to anodes through electrodeposition [32, 76], catalysis of solvent reduction [77] and formation of inorganic layers in SEIs [78], all of which can impede Li+ intercalation and reduce capacity [79]. Here, Mn3+ disproportion into Mn2+ and Mn4+ was proposed to explain Mn dissolution [26, 75] as previously reported in Li–Mn–O spinel structures [80, 81] and various electrolytes have been utilized to form protective surface films containing carbon on electrodes to isolate surfaces from electrolytes [33].

$$ {\text{LiPF}}_{6} \to {\text{LiF}} + {\text{PF}}_{5} $$


$$ {\text{LiPF}}_{6} + 2{\text{e}}^{ - } + 2{\text{Li}}^{ + } \to {\text{LiF}} + {\text{Li}}_{x} {\text{PF}}_{y} $$


Self-redox Reaction

At highly delithiated states, the valence of transition metals increases, and because the low-spin Co3+/4+:t2g band overlaps with the 2p band of O2− (Fig. 3b), electron transfer from O2− to Co3+/4+ can occur and reduce oxidized transition metal ions, thus releasing O2 into battery cells [82]. Oxygen release based on this also occurs at the first charging period in which O2− reduces transition metals oxidized to 4+ (similar effects on Ni4+ also exist and dominate oxygen release in Ni-rich NMC-based cathodes) [14, 42].

Parasitic Reaction: Effects of Ni4+

Highly delithiated (highly charged) states tend to generate large amounts of Ni4+ that can react with electrolytes. (No reaction schemes have been proposed.) This side reaction can significantly thicken cathode–electrolyte interfaces (CEIs) and reduce the number of available Li+ [14, 83], thus increasing impedance [13]. Researchers have also investigated gas evolution (CO2) from NMC532 cathodes using DEMS [84], and isothermal calorimetry (IMC) results have shown that at highly delithiated states above 4.2 V, the further removal of Li from NMC811 cathodes can result in decreased entropy and endothermic heat flow [7] in which rapidly increasing heat flow indicates that the highly delithiated cathode is very reactive with the electrolyte and can therefore cause capacity fading. In addition, researchers also found that remaining at a highly delithiated state can cause significant electrolyte oxidation as well as other side reactions [9].

Parasitic Reaction: CO2generation

Large Ni4+ portions at higher voltages can cause increased CO2 evolution at increased upper cutoff voltages [85] in which cathodes with higher Ni content suffer more from parasitic reactions with Ni4+ that are a source of emitted CO2. In addition, because highly reactive oxygen can be produced through O2− from self-redox reactions, surface reconstructions and highly delithiated cathodes (ca. > 80%) [86], CO and CO2 can also be traced back to the reaction between O2 and alkyl carbonate electrolytes (Eq. 4) [24, 25]. Alternatively, sweeping to higher voltages can also cause the electrochemical decomposition of electrolytes to produce CO2 (Eq. 5, Fig. 3c) [23, 24, 87]. Furthermore, Gasteiger et al. [88] have experimentally shown three other mechanisms that may contribute to CO2 emission, including solvent hydrolysis (if trace amounts of OH− exist), electrolyte impurity oxidation and HF reacting with lithium carbonate as a surface impurity. Moreover, the decomposition of lithium carbonate has also proven to be a source of CO2 [89, 90]. Researchers have also suggested that NMC-based cathodes with Ni-rich surfaces behave differently than other cathodes with different portions of Ni [7]. For example, the degradation of NMC811 cells was studied by using different electrolyte additives including VC and PES211 to suppress side reactions and it was revealed that NMC442/graphite cells and NMC111/graphite cells showed better performances with PES211, whereas NMC811/graphite cells aged slower with VC (with slower capacity fading and impedance growth) [79]. Here, PES was also reported to be a viable additive for the suppression of gas evolution [91] and impedance growth during cycling to higher voltages (above 4.3 V), and rocksalt surface layers formed with the use of VC but not with PES211, concluding that electrolyte additives can significantly affect the rate of parasitic reactions and that parasitic reactions were the main reason for capacity fading [7, 92].

$$ {\text{EC}} + 2{\text{O}}_{2} \to 2{\text{CO}}_{2} + {\text{CO}} + 2{\text{H}}_{2} {\text{O}} $$


$$ {\text{EC}} \to {\text{CO}}/{\text{CO}}_{2} + {\text{R-H}}^{ + } + {\text{e}}^{ - } $$


Other Issues

Initial Capacity Loss

Initial capacity loss is strongly linked to the loss of available Li and the growth of impedance in anodes due to SEI formation. Here, researchers have reported that NMC-based cathodes cannot return to their fully lithiated states after initial cycling [26] in which separate mechanisms have been proposed, including parasitic reactions consuming Li+ [93] and the sluggishness of Li+ diffusion into the few vacancies of the Li-layer [94] (circumstantially proved by Gasteiger et al. [26]). In addition, researchers have also suggested that initial capacity loss is related to the C-rate because it is kinetically influenced. Furthermore, partial losses of available cathodic Li occur to compensate for losses of anodic Li+ to SEIs [26], and other losses of Li in the cycling period can be attributed to parasitic reactions including the immobilization of Li in the SEI and enlarged polarization, which become significant at higher cutoff potentials [26, 95].

Cracking of Secondary Particles

Reduced performances due to the cracking of spherical secondary particles have been extensively reported and are of significant concern [96]. Here, the expansion and contraction of the c-axis during repeated cycling at above 4.2 V can lead to microstrains on particles in cathodes, leading to the generation of microcracks in the core of primary particles with some initially generated cracks being able to close in subsequent cycles [68]. These cracks can cause poor connection between particles and micropores and even the cracking of secondary particles [13], all of which contribute to degradation at the microscale. To address this, Lim et al. [27] have used first-principle calculations to study these cracks generated by gaps between primary particles and found that they were a result of anisotropic and average contractions. In addition, researchers reported that these cracked particles possessed enlarged surface areas and therefore increased the possibility of parasitic reactions [13]. Moreover, these microcracks were also found to partially originate from the gas evolution of NMC particles [13] (Fig. 3d), and recent studies on coated NMC (76 14 10) suggested that the infiltration of liquid electrolytes into the gap between primary particles can also cause cracking [97]. Alternatively, various NMCs have also been found with no cracking by using SEM [7, 20]. Here, Dahn et al. [20] suggested that the shell of core–shell structured NMC cathodes (a Mn-rich shell and a Ni-rich core) was resistant to cracking due to the absence of compressive stress on the core and tensile stress on the shell. Based on all of this, numerous mitigation methods including carbon matrixes and gradient layers have been proposed to resolve cracking at the microscale (further details in Sect. 4.1).

The Effects of Temperature

Studies have shown that cell degradation can accelerate at elevated temperatures above 30 °C and that at low C-rate cycling, cell degradation is mainly affected by the time in which cells are placed under high temperatures and not by cycling [98]. Long-term operation of LIB cells can produce large amounts of heat. And although this released heat may even be useful in cold weather conditions, heat emissions are a major issue of Ni-rich NMC-based LIBs due to thermal instability [33] in which thermal runaways can occur if NMC-based cathodes react with LiPF6 from the electrolyte and limit the commercial and practical use of these battery packs. To improve thermal stability, coatings such as SiO2 [62], TiO2 [99] and Li2ZrO3 [100] have been chosen to prevent contact between cathodes and electrolytes, and an ALD-coated NMC532 with Al2O3 has even been reported [101] to resolve the issue of uneven coatings and prevent the failure of cathode separation from electrolytes.

Graphite Anode

Similar to the cracking of NMC secondary particles, cracking can also occur in graphite anodes due to Li intercalation/deintercalation, which can expand and contract the distance between graphite layers [102]. In addition, recent research indicated that unit cell expansion (13.2% in total) during Li intercalation is neither continuous nor linear, but staged (Fig. 4a–e) [103]. Here, Li intercalation from C6 (Li-free graphite) to LiC24 (Stage 2L) showed no in-plane ordering of intercalated Li, whereas from 2L to LiC12 (Stage 2), the ordering of Li appeared and the lattice can be assigned to the space group P6/mmm. Furthermore, in the transition from Stage 2 to LiC6 (Stage 1), total volume change suddenly increases as the amount of x exceeds 0.6 in LixC6 (where Stage 1 appears as Li presents in every interlayer). These results showed that optimal composition was between 20 and 80% charged in which Stage 2 stably exists or extends into.

Fig. 4

Crystal structures of a graphite (space group P63/mmc) and two major Li intercalation compounds, b LiC12 and c LiC6 (space group P6/mmm). Different stacking of graphene layers with d the AB sequence in graphite and e the AA sequence in LiC12 and LiC6 (Reprinted with permission from Ref. [103]. Copyright © 2018, American Chemical Society). f Schematic of SEI formation and degradation on graphite anodes (Reprinted with permission from Ref. [31]. Copyright © 2005, Elsevier B.V. All rights reserved.)

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Furthermore, “dead lithium,” which is Li-ions electrochemically plated onto the surface of graphite anodes, is the principal reason for capacity fading, especially at final stages [104]. This phenomenon occurs on graphite surfaces that possess limited mass transport, slow diffusion within the graphite or low charge transfer speeds [105] and is strongly promoted by low temperatures/large polarizations (including high electrolyte resistances induced by high temperatures) during lithiation as well as the properties of the graphite electrode (e.g., morphology [102], loading). And as Li+ intercalation is replaced by Li plating, issues such as dendrite growth become more pronounced, leading to increased impedance and more importantly, potential safety hazards due to internal short circuiting. To avoid this, the limits of charging currents and the state of charge need to be established to maintain polarization in optimal ranges [106]. However, large polarizations also need to be adopted to minimize charging durations to meet industrial application demands. Here, researchers reported that the reduction of graphite thicknesses and improvements in porosity can decrease charge transfer impedances [105, 107].


Electrolytes used in LIBs contain two main components including Li salt and organic solvents. Here, the use of additives as film-forming agents can allow for stable SEIs and, in general, the concentration of Li salt is maintained at 0.8–1.3 M to optimize electrolyte viscosity and ion conductivity. As compared with cathodes, the potential of LIB anodes is low and can trigger electrolyte reduction (SEI formation), which can cause Li loss during initial charging (Fig. 4f) [31]. However, because SEIs are electronically insulating, further reductions of electrolytes do not occur in which formed SEIs are usually 10–50 nm thick and are conductive to Li+ ions. Researchers have extensively studied the constituents of passivating SEIs on graphite anodes and have reported that lithium ethylene dicarbonate (LEDC, Eq. 6, Fig. 5a–d) [108] was the main component of SEIs formed with ethylene carbonate (EC) solvent [109] apart from electronically insulating LiF and gaseous products (H2 and C2H4). Propylene carbonate (PC) as a solvent has also been well investigated due to its thermal stability and wide voltage windows. Furthermore, DFT calculations have shown that stable SEI films can be produced by oligomers of SEI film components (SFCs) [108], which are first created through solvent decomposition. LiF is also a product of almost all common Li salts such as LiPF6 (the only commercially available salt), LiBF4, LiTFSi, LiTSi, etc., in which researchers reported that the amount of LiF produced from LiBF4 was large and would result in the formation of a grainy film instead of a smooth film.

Fig. 5

DFT-MD snapshots of electrode/electrolyte interphases with Li2EDC: a adhesion structure of one Li2EDC monomer on a graphite electrode, b dissolution structure of one Li2EDC molecule in EC solvent, c adhesion structure of 12 Li2EDC molecule aggregates on a graphite anode, d dissolution structure of 12 Li2EDC molecule aggregates in EC solvent (Reprinted with permission from Ref. [108]. Copyright © The Author(s) 2015. Published by ECS). e Proposed mechanism for the continuous decomposition of SEI and electrolytes as monitored by C2H4 evolution for a preformed electrode with a Mn2+-containing electrolyte: (1) absorption of Mn2+ ions into the SEI; (2) reduction of Mn2+ ions in the SEI and deintercalation of Li+ from graphite; (3) re-oxidation of Mn0 to Mn2+; (4) recurrent electrolyte reduction; (5) the catalytic cycle of electrolyte decomposition (Reprinted with permission from Ref. [38]. Copyright © The Author(s) 2018. Published by ECS)

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$$ 2{\text{EC}} + 2{\text{Li}}^{ + } + 2{\text{e}}^{ - } \to {\text{Li}}_{2} {\text{EDC}} + {\text{C}}_{2} {\text{H}}_{4} $$


As aging continues during cycling, researchers believe that the inorganic portions of SEIs also increase as LiF (can be used as an indicator of SEI maturation) is produced through LiPF6 reacting with lithium carbonates (Eqs. 7, 8) [37].

$$ 2{\text{RCO}}_{3} {\text{Li}} + {\text{LiPF}}_{6} \to 2{\text{LiF}} + 2{\text{RF}} + {\text{LiPO}}_{2} {\text{F}}_{2} + 2{\text{CO}}_{2} $$


$$ {\text{RCO}}_{3} {\text{Li}} + {\text{LiPF}}_{6} \to 2{\text{LiF}} + {\text{RF}} + {\text{POF}}_{3} + {\text{CO}}_{2} $$


After SEI formation, the resistance of anodes/SEIs toward divalent transition metal ions such as Ni2+, Co2+, Cu2+ and Mn2+ (relatively more impactive) is significantly enhanced, preventing the reduction of EC solvent by these ions. However, formed SEIs remain targets to many reductive substances in which reductive decomposition can degrade SEIs (Li2CO3 and C2H4 are produced as a result) and expose thermodynamically unstable graphite anodes to reactive environments [38]. Here, if the exposure of graphite anodes occurs, a series of damage can occur on the graphite anode, including the intercalation of solvated Li+, which leads to larger sizes and the potential risk of the breaking of the graphite layer. In addition, Mn2+ has been reported to be able to penetrate SEIs and reside in the inner area between the SEI and the anode through the exchange of positions with Li+ ions located in the SEI [110]. Furthermore, the recurrent generation of C2H4 suggests that Mn2+ is a destructive substance to both solvents and graphite anodes in which accumulated Mn2+ near graphite can function as a catalyst, suggesting that it is reduced at the surface of the anode but can reduce EC afterward, leading to thickened SEI films and the immobilization of active Li (Fig. 5e) [110]. In addition, the reduction of Mn2+ can cause the gradual transfer of active Li+ to SEIs or the conversion of active Li+ into compounds deposited onto graphite anodes, both of which lower storage capacity [38]. Researchers have also proposed another mechanism of Li+ that suggests that even monolayer Mn2+ can impede the movement of Li+ through the graphite and the SEI [111]. As for PF5, it is produced from the decomposition of various electrolytes and can trigger the open-ring polymerization of EC, which further increases SEI decomposition along with the production of HF [112]. Moreover, the damage/repair process caused by gaseous PF5 consumes active Li-ions until Li-ions are depleted in the system [113]. Researchers have reported, however, that this degradation can be mitigated through the addition of Lewis-basic additives. The presence of H2O in electrolytes can also reduce the production of C2H4 (usually used as an indicator of EC reduction) and promote H2 production. Here, the existence of SEIs can reduce the reduction of trace amounts of water that usually generate H2 and CO2 [114].

An automotive scale NMC pouch cell was also analyzed by Dahn group using ultra-high precision cycling technologies, and it was revealed that the performance of graphite anodes was decisive and that corresponding capacity fading (proportional to t1/2) was related to SEI thickness and was inversely proportional to the rate of SEI growth, suggesting that the gradual lowering of SEI thickening rates can lead to stable SEIs [115, 116]. Here, alternative relationships proposed for SEI growth need to take stability as a factor into consideration [117].


Overall, the degradation of Ni-rich NMC-based cathodes originates from the insertion/desertion of Li+ during charge/discharge in which the expansion and contraction of crystal lattices (primary particles) gradually generates tension, leading to the cracking of secondary particles (Fig. 6). Electrolyte infiltration as well as oxygen release can further deteriorate this issue. On an atomic scale, Ni4+ can accumulate during charging because delithiation can elevate the valence state of Ni together with high-valence Mn and Co ions. And due to smaller radii, easier diffusion can lead to surface reconstruction (from rhombohedral to rocksalt), which can further inhibit Li+ insertion/desertion. Moreover, increased voltage levels of high-valence ions can trigger side reactions such as the oxidization/decomposition of electrolytes and the further dissolution of NMC materials. Electrochemical potentials of electron-depleted Co and Ni can also fall below the fermi level of O2−, leading to the transfer of electrons from O2− to transition metal ions and the release of O2 into systems, which can also cause electrolyte decomposition. Throughout the operation of Ni-rich NMC-based LIBs, four types of unwanted substances exist on the surface of cathodic particles, including other phases, surface impurities (carbonates and hydroxide), rocksalt structures and surface films. To sum up, aging effects on both micrometer and atomic scales can result in two common outcomes in which the first is the chemical (through chemical reactions) and mechanical (through the breakage of crystallites) damage/loss of active materials and the second is the production of detrimental substances on electrode particle surfaces, including surface films (both SEIs and CEIs), surface impurities and surface layers. And although many differences in the performance of various Ni-rich NMC-based cathodes have been discovered, these remain unexplained. Therefore, future studies should focus on the interpretation of different aging behaviors in NMCs with or without mitigation methods.

Fig. 6

Summary of degradation mechanisms of Ni-rich NMC-based cathodes on the micrometer and atomic scale

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As for graphite anodes, degradation primarily occurs due to the plating of Li dendrites and the cracking of layered graphitic structures. In addition, SEIs tend to thicken during cell operation as EC solvent is reduced electrochemically. Furthermore, transition metals, especially Mn2+ from dissolved NMC, can cause significant capacity loss due to massive damage to SEIs. A comprehensive list of identified degradation mechanisms for Ni-rich NMC-based cathodes, graphite anodes, SEIs and CEIs is presented in Table 1. Here, the degradation mechanisms of SEIs are categorized based on anodes and the degradation mechanisms of CEIs are categorized based on cathodes because certain degradation mechanisms may involve both electrode materials and interfaces.

Table 1 List of identified degradation mechanisms of Ni-rich NMC/graphite LIBs

Full size table