Addressing the fundamental issues in Ni-rich cathodes: degradation mechanisms and mitigation strategies

Abstract

The rapid growth of electric vehicles (EVs) is driving an urgent demand for lithium-ion batteries (LIBs) with higher specific energy, longer life, and uncompromised safety. Ni-rich layered oxides (LiNi_xC_oyMn_(1−x−y)O₂, x ≥ 0.8) have emerged as leading cathode materials for next-generation LIBs, owing to their high capacity and energy density. Further increasing Ni content is essential for improved performance and cost reduction. However, it also introduces new obstacles, necessitating thoughtful design of cathode composition, morphology, and microstructure, as well as the development of electrolyte formulations. In this review, we discuss the multiple failure mechanisms of Ni-rich cathodes in terms of two major aspects: structural degradation and gas release. We elucidate the key factors contributing to chemical, crystallographic, and microstructural degradation in Ni-rich cathodes, and summarize the various origins of gas evolution associated with these materials. Another key theme of this review is the modification of Ni-rich cathodes to address the practical hurdles that limit their use in long-range and high-safety EVs. Accordingly, we present a comprehensive overview of the latest Ni-rich cathode modification strategies for next-generation EV platforms.

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Ziqi Liu

Ziqi Liu is currently pursuing a PhD at the School of Materials Science and Engineering, Sichuan University, under the supervision of Professor Yun Zhang. Her research focuses on the field of cathode materials for lithium-ion batteries.

Yiming Zhang

Yiming Zhang is currently a PhD candidate at the School of Materials Science and Engineering, Sichuan University, under the supervision of Professor Jiagang Wu. His research focuses on the recycling and comprehensive utilization of lithium-ion battery cathode materials.

Yong Chen

Yong Chen is currently pursuing his PhD at the Centre for Clean Energy Technology, Faculty of Science, University of Technology Sydney, under the supervision of Prof. Guoxiu Wang. His research focuses on solid-state batteries and electrode materials. He has published more than 10 papers as the first author or corresponding author in high-impact international SCI journals, including Energy Storage Materials, Advanced Functional Materials, Advanced Materials (2), Angewandte Chemie International Edition (2), Advanced Science, Energy & Environmental Science (2), Journal of Energy Chemistry, and Journal of Power Sources.

Shuaicheng Jiang

Shuaicheng Jiang, PhD, graduated from the College of Materials Science and Engineering of Nanjing Forestry University in June 2023 and has been working in Beijing Forestry University since October 2023. His current research interests focus on biomass-based nanocomposite materials with sustainability, renewability and biocompatibility, including elastomers, membranes, and adhesives, and their advanced applications in the fields of environmental protection, green energy, flexible electronics, etc.

Guoxiu Wang

Prof. Guoxiu Wang is a Distinguished Professor at the University of Technology Sydney (UTS) and Director of the Centre for Clean Energy Technology. He is a Fellow of the European Academy of Sciences, the Australian Academy of Science, the Australian Academy of Technological Sciences and Engineering, the International Society of Electrochemistry, and the Royal Society of Chemistry (UK). Prof. Wang has long been devoted to the fields of materials chemistry, electrochemistry, energy storage and conversion, battery technology, and catalysis. He currently serves as Associate Editor of Electrochemical Energy Reviews and Energy Storage Materials. To date, he has published over 700 scientific papers with more than 94 000 citations and an h-index of 167 (Google Scholar), and has been recognized as a Highly Cited Researcher (Materials Science and Chemistry) by Web of Science/Clarivate Analytics.

Meng Yao

Meng Yao is an Associate Researcher at the School of Materials Science and Engineering, Sichuan University. His research focuses on solid-state electrochemical energy storage and lithium extraction from salt lakes. To date, he has published over 20 academic papers as the first/corresponding author in high-impact journals including Energy Environ. Sci., Angew. Chem. Int. Ed., Adv. Energy Mater., Adv. Funct. Mater., Energy Stor. Mater., and Small. He has led or participated as a principal investigator in three National Natural Science Foundation projects, two provincial/ministerial-level research initiatives, and one industry-sponsored research project.

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Broader context

Lithium-ion batteries (LIBs), serving as the cornerstone of contemporary electrochemical energy storage systems, have gained widespread adoption across consumer electronics and electric vehicles (EVs) owing to their superior energy density, extended cycle life, and minimal self-discharge characteristics. Cathode materials play a decisive role in determining key battery performance metrics. Their technological evolution is connected to the realization of global energy transition objectives and carbon neutrality targets. Innovations in cathode materials will continue to push the boundaries of battery performance. From a sustainable development perspective, the development of cathode materials must balance resource availability, environmental friendliness, and ease of recycling. Among various cathode materials, Ni-rich materials have emerged as a promising LIB cathode material due to their high energy density, low cost, and strong industrial compatibility. However, the high nickel content leads to severe structural degradation and vigorous gas evolution, limiting their practical application. This paper aims to comprehensively elucidate the reaction mechanisms underlying structural degradation and gas release in Ni-rich cathode materials and summarize recent solutions addressing these two issues. We hope to provide valuable references for future research on Ni-rich cathode materials from a practical application perspective.

1. Introduction

With the continuous transformation of the global energy system, conventional fossil fuels are progressively being displaced by cleaner, low-carbon alternatives to meet growing energy demands.^1–4 The development of efficient energy conversion and storage devices is a critical approach to addressing the intermittency and uncertainty shortcomings of clean energy itself.^5,6 Among numerous energy storage technologies, lithium-ion batteries (LIBs) have been widely adopted in various essential domains, due to their high energy density, long cycle life, and lack of memory effect.⁷ Especially, LIBs are also considered the battery of choice for powering the next generation of hybrid electric vehicles and plug-in hybrid electric vehicles and the global energy storage capacity of electric vehicles (EVs) has grown rapidly in recent years,⁸ with projections estimating approximately 240 million units by 2030 (Fig. 1a).⁹ This surge in demand presents unprecedented opportunities for the development of LIBs. Therefore, to meet the growing demands for longer driving range and enhanced safety in EVs, the development of high-performance LIBs has become a key focus for both academia and industry.

Fig. 1 (a) Projected global electric car stock compared with original equipment manufacturer (OEM) targets (2020–2025).⁹ Copyright 2019, International Energy Agency. (b) A breakdown of production/material costs of a typical LIB cell.¹⁰ Copyright 2018, Springer Nature. (c) Cycle life and energy density of a range of fully or partially commercialized LIB cathodes. (d) History of LIB cathodes.

As a key component of LIBs, the cathode materials are the main cost driver of the battery (Fig. 1b), and are the main factor determining the energy density and safety of the battery.^10,11 The commercial LIB cathode materials mainly include lithium manganate (LiMn₂O₄, LMO),^12,13 lithium iron phosphate (LiFePO₄, LFP),^14,15 lithium cobalt oxide (LiCoO₂, LCO),^16,17 and ternary cathode materials (LiNi_xCo_yMn/Al_zO₂, x + y + z = 1, NCM/NCA) with different Ni contents.^18–24 Considering the actual needs of power batteries and the current technical sophistication, the Ni-rich cathode (LiNi_xCo_yMn/Al_zO₂, x + y + z = 1, x ≥ 0.8) is becoming one of the most promising cathode materials for the next generation of power batteries due to its high theoretical discharge capacity (275 mAh g⁻¹), high operating voltage (3.8 V), and low cost.^25–27 To further increase the capacity and energy density, ultrahigh Ni cathodes with a Ni content of more than 90% are also promising next-generation cathodes, especially under high voltage (HV Ni-rich cathodes).^28–31 In addition, considering the high price of Co, Co-free materials are also potential options for future LIB cathodes (Fig. 1d).^32–40 As shown in Fig. 1c, Ni-rich cathodes and ultrahigh Ni-rich cathodes exhibit satisfactory cycle stability and excellent energy density, and are expected to be applied on a large scale in EVs. On the other hand, Co-free cathodes, such as lithium-rich NCM (LR-NCM) and LiNi_0.5Mn_1.5O₄ (LMNO), are only considered as emerging cathode materials due to their shorter cycle life or lower energy density.^29,41–44

Although Ni-rich cathodes promise compelling energy gains, their large-scale deployment demands stringent mitigation of the diverse degradation phenomena that escalate with higher Ni content. Previous studies typically summarized the problems of Ni-rich cathodes as independent points, including cation mixing, excessive residual lithium contents on the surface, phase transition and microcrack generation, poor safety performance, etc.^45–48 However, after comprehensive analysis in this review, it is found that the problems arising from Ni-rich cathodes can be divided into two major categories: structural degradation problems, including (1) generation of residual lithium compounds (RLCs) on the surface,^49–52 (2) interfacial side reactions,^53–57 (3) surface phase transition,^58–62 (4) microcracks and particle cracking,^63–66 and (5) cation mixing,^60,67–70 and gas production issues, including gases such as CO, CO₂, and O₂.^62,71,72 In order to alleviate the structural decay and gas release existing in Ni-rich cathodes, the researchers adopted various strategies. Surface coating could prevent electrolyte erosion and oxygen release by forming a protective layer on the surface of the cathodes.^73–76 Element doping is also an important means to improve structural stability and inhibit gas production. Doping of low valence metal ions (e.g., Na⁺, Mg²⁺) could inhibit the collapse of layered structures in the form of “pillars”.^77–81 Some high valence transition metal ions (e.g., Ta⁵⁺, Nb⁵⁺, W⁶⁺, Mo⁶⁺, etc.) could refine the primary grains, ease stress accumulation, and enhance the bonding with oxygen atoms, thus improving the lattice oxygen stability.^{28,46,78,81–86} Besides, optimization of the electrolyte by adding specific additives (e.g., fluorine-rich additives, phosphate additives, and nitrile-based additives)^87–91 and rational structural design of the Ni-rich cathodes^64,92–95 could also effectively inhibit the release of gases and stabilize the structure of the cathode materials.

Focusing on the structural stability and gas emission behavior of Ni-rich layered cathodes, this review analyses the detailed mechanisms of structural damage as well as the gas production mechanisms and modification progress of Ni-rich cathodes. Firstly, we systematically summarize the causes, processes and effects of various types of structural damage, including chemical composition changes and lattice structure transformation on the surface/interface of Ni-rich cathodes, as well as microcracks and cation mixing problems throughout the whole Ni-rich cathodes. Then, multiple mechanisms of gas emission from Ni-rich cathodes are summarized in detail, including the decomposition of RLCs and electrolytes as well as the release of lattice oxygen and the effects of gas emission. Finally, this paper also summarizes strategies to suppress the structural damage and gas emission behavior of Ni-rich cathodes based on the failure mechanism. Through surface modification, element doping, electrolyte design and structural engineering, the safety and cycle life of the battery can be effectively improved. By understanding the failure mechanism of Ni-rich cathodes coupled with each other, researchers can adopt richer and targeted modification strategies to provide scientific basis and technical support for the commercial application of Ni-rich cathode materials.

2. Structural damage mechanisms of Ni-rich cathodes

2.1. Degradation of the interface/surface structure

2.1.1. Residual lithium compounds. The formation of RLCs on the surface is one of the common problems in Ni-rich cathode materials.⁹⁶ The RLCs may originate from two aspects. First, during the preparation of layered Ni-rich cathode materials, due to the volatilization and loss of Li during high-temperature sintering, excess Li favors the formation of highly ordered Ni-rich compounds.⁹⁷ Unreacted Li mainly exists as Li₂O on the surface of the material after the calcination is completed. After exposure of the material to air, Li₂O comes in contact with CO₂ and H₂O and is converted LiOH, LiHCO₃, and Li₂CO₃; the reaction is described by eqn (1)–(3)⁹⁸ in Fig. 2a. Secondly, during the cooling process and storage, the surface layer undergoes remodeling, and Li⁺ in the lattice would react with H₂O and CO₂ in the air to form a large number of RLCs. At present, there are three possible mechanisms. There are also multiple perspectives on the mechanism by which surface remodeling leads to the formation of RLCs. The first view is that the formation of surface RLCs is related to surface lattice oxygen. Liu et al.^99,100 concluded that Ni-rich cathodes spontaneously reduce Ni³⁺ to Ni²⁺ during storage and proposed a surface reaction mechanism based on the Ni³⁺/Ni²⁺ transition and the evolution of reactive oxygen species. After the spontaneous reduction of Ni³⁺ to Ni²⁺, the lattice oxygen anion O²⁻ is oxidized to the reactive oxygen anion O⁻. Then, the active oxygen anion O⁻ undergoes a disproportionation reaction to produce the active oxygen anion O²⁻ and the neutral oxygen atom O; the active oxygen anion on the surface combines with CO₂ and H₂O in the air to form CO₃²⁻ and OH⁻, and the neutral oxygen atom O combines with the active oxygen anion O⁻ to form the surface-active oxygen anions. Alternatively, the neutral oxygen atom O may reorganize to O₂, and finally, CO₃²⁻ and OH⁻ react with surface Li⁺ to form Li₂CO₃ and LiOH on the surface of the Ni-rich cathodes. The reaction process is described by eqn (4)–(8) in Fig. 2a. This potential mechanism is further supported by the study of Abraham et al.,¹⁰¹ which identified a thin NiO-like layer near the surface region between the adsorbed species layer (hydroxyl, carbonate, and bicarbonate) and the bulk phase. Kaufman et al.¹⁰² further confirmed through isotope labeling that reactive oxygen species released from the lattice may be triggered by the oxidation of RLCs.

Fig. 2 (a) Formation process of RLCs on the surface of Ni-rich cathodes. (b) Sources of surface RLCs and their multiple formation mechanisms.¹⁰⁷ Copyright 2020, John Wiley and Sons. (c) Molecular mechanism of PVDF degradation.¹⁰⁸ Copyright 2020, American Chemical Society. (d) Schematic illustration of slit extrusion coating operation.¹⁰⁹ Copyright 2020, Elsevier. (e) Schematic diagram of electrode surface defects.

The second viewpoint, on the other hand, suggests that H₂O in the air could interact with CO₂ to form Li₂CO₃ on the surface of the Ni-rich cathodes. This viewpoint assumes that Li₂CO₃ on the surface of the Ni-rich cathodes is generated in the presence of H₂O at room temperature. H₂O has a catalytic effect and could react with CO₂ in the air to form H₂CO₃, which leads to material surface corrosion due to the acidic nature of the surface water. As a result, Li⁺ is easily leached from the material surface lattice and reacts with CO₃²⁻ to form LiOH and Li₂CO₃. The reaction process is described by eqn (9)–(12) in Fig. 2a.^103–105

The third viewpoint suggests that the formation of RLCs may be due to the exchange of Li⁺ and H⁺ protons near the surface of the Ni-rich cathodes. During storage, Li⁺ migrates to the surface of the material and reacts with H₂O and CO₂ in the air to form RLCs, and the released H⁺ occupies the Li⁺ position. The reaction process is described by eqn (13)–(15)¹⁰⁶ in Fig. 2a.

With the formation of RLCs, the Li⁺ in the cathode surface structure would continue to be lost, which would cause the degradation of the electrochemical performance of the battery. Moreover, reactive RLC tends to undergo parasitic reactions with the electrolyte, resulting in gas generation and interfacial instability. RLCs (LiOH, Li₂CO₃) exhibit alkaline properties. While the widely used polyvinylidene fluoride (PVDF) binder demonstrates strong resistance to acids, oxidants, and organic solvents, it undergoes significant structural degradation in alkaline environments, leading to the phenomenon of slurry gelation (Fig. 2b).^{107,108,110–112} The PVDF molecular chain contains carbon atoms that are not fluorinated. Under the strong electron-withdrawing inductive effect of fluorine atoms, the hydrogen atoms on the carbon atoms become positively charged. Consequently, they readily undergo attack by strong nucleophiles (OH⁻), leading to a hydrogen fluoride elimination reaction that generates a conjugated polyene structure. This polyene structure, composed of conjugated double bonds, makes it susceptible to subsequent nucleophilic attack by OH⁻ ions. The generated double bonds induce cross-linking reactions between PVDF molecular chains, ultimately forming a gel as shown in Fig. 2c.^108,113 The high regularity of the PVDF chain leads to a chain reaction after dehydrofluorination, accelerating both the dehydrofluorination and gelation processes.¹¹⁴

In actual production, the gelation of cathode slurry caused by RLCs leads to a sharp increase in its viscosity, causing the slurry to lose its fluidity. When preparing electrode sheets using the currently mainstream slit extrusion coating apparatus as shown in Fig. 2d,¹⁰⁹ this may result in the slurry becoming impossible to pump through pipelines and may cause uneven flow resistance or even complete blockage during the coating process. This unpredictable processability makes stable, continuous slurry delivery and coating impossible. Furthermore, partially gelled slurry leads to gel particles forming on the coated electrode surface, resulting in abnormal porosity and uneven thickness (Fig. 2e).¹¹⁵ These macro defects directly cause uneven current distribution within the battery, accelerate performance degradation, and lead to poor consistency between individual cells.¹¹⁶ Furthermore, slurry gelation may also force production lines to frequently clean slurry delivery and coating systems, significantly reducing production efficiency and substantially increasing manufacturing costs due to slurry waste and lost production capacity.

However, recent studies have shown that Li₂CO₃ is the most thermodynamically stable RLC, and its solubility in organic solvents is much lower than that of LiOH, which could prevent OH⁻ from attacking the PVDF. In addition, Li₂CO₃ would react with HF to form a F-rich cathode electrolyte interface (CEI). Therefore, an appropriate amount of Li₂CO₃ can effectively enhance the air stability and cycling performance of Ni-rich cathodes.¹¹⁷

2.1.2. Side reactions at the interface. Another structural degradation phenomenon in Ni-rich cathodes is the chemical structural degradation, which leads to the formation of the CEI at the electrode/electrolyte interface. This is due to a series of interfacial side reactions such as electrolyte decomposition and dissolution of transition metals (TMs) during the cycling process.

Interfacial side reactions between the cathode and the electrolyte are responsible for the formation and evolution of the CEI on the surface of Ni-rich cathodes. The formation of the CEI is closely related to the oxidation of the electrolyte. Primary CEI formation occurs when the cathode material is in contact with the electrolyte. It has been shown that LiNiO₂ leads to the oxidation of the electrolyte and thus the formation of the CEI due to its strong nucleophilic oxygen atoms and Lewis base properties.¹¹⁸ In addition, according to previous studies, electrochemical oxidation of the electrolyte occurs when the lowest unoccupied molecular orbital (LUMO) energy level of the cathode active material is lower than the highest occupied molecular orbital (HOMO) energy level of the electrolyte. Since the energy of the antibonding hybridization orbitals of the TM 3d-O 2p of the Ni-rich cathode in the charging state corresponds to the LUMO of the holes and the concentration of holes inside the Ni-rich cathode increases, the LUMO energy level of the cathode decreases continuously, gradually approaching the HOMO level of the electrolyte, which may lead to the oxidation of the electrolyte as shown in Fig. 3a.¹¹⁹ In this case, the CEI of the Ni-rich cathode consists of Li₂CO₃ and alkyl carbonate oxidized from the electrolyte solvent and Li_xPO_yF_z oxidized from the electrolyte salt, and it continues to grow during the charge/discharge cycle. Besides, in the deep delithiation state, the presence of highly oxidized and unstable Ni⁴⁺ further promotes the oxidation of the electrolyte. As the Fermi energy level of the material decreases, the O²⁻ ions lose electrons to become highly reactive O or O₂, which also exacerbates electrolyte depletion and the formation of a thick and inhomogeneous CEI.

Fig. 3 (a) Schematic of the oxidation decomposition of the electrolyte on the surface of cathodes. (b) Transition metal dissolution mechanisms in Ni-rich cathodes. (c) Schematic illustration of the origin and component of the CEI layer. (d) Architecture comparison of CEI layers formed at low/high voltage.¹²³ Copyright 2023, American Chemical Society.

In addition, the dissolution of TMs is a typical interfacial side reaction and promotes the formation of a CEI. After prolonged cycling, the TMs on the surface of the Ni-rich cathode will dissolve out of the lattice. This phenomenon is mainly affected by two factors: on the one hand, the cation mixing and oxygen evolution reactions generated during the cycling process lead to the formation of low-valent transition metal oxides, which have high solubility in the electrolyte.⁶² On the other hand, the electrolyte solvent contains traces of water, and these traces of water would react with LiPF₆ to form HF, which dissociates into H⁺ and F⁻ and frequently attacks the O–Li and O–TM (TM: Ni, Co, Mn) bonds to break them. The fractured O²⁻ ions combine with H⁺ to form H₂O, and the TM ions are detached from the crystals, which eventually dissolve in the electrolyte or combine with F⁻ to form TMF₂, which is deposited on the surface of the Ni-rich cathode to form a CEI. In the process, H₂O is generated, which further promotes the solvation of LiPF₆, forming a self-driven cyclic reaction as shown in Fig. 3b.¹²⁰ However, in the case of NCM type cathodes, the direct loss of the active substance plays only a small role in the capacity decay. The secondary effects caused by the migration and deposition of dissolved TMs on the cathode surface are the main cause of the degradation of Ni-rich cathode performance. As a result, the active lithium on the anode side would be consumed and produce uneven deposition, resulting in lithium dendrites, which would lead to poor battery cycling performance as well as short circuits and thermal runaway, thus posing safety risks.^121,122

The above interfacial side reactions yield abundant organic fragments and inorganic by-products, ultimately forming a composite CEI in which an inorganic-rich inner solid-electrolyte layer (SEL) and an organic-rich outer quasi-solid-electrolyte layer (qSEL) merge without a distinct boundary. The specific components of the CEI are shown in Fig. 3c.¹²³ In particular, the SEL consists mainly of oxides, fluorides, fluorophosphates, carbonates, and other dense inorganic materials derived from the complete decomposition of salts and solvents, while the qSEL consists mainly of incompletely decomposed organic products such as polyolefins, semi carbonates and polymers.^124,125 In fact, the specific structure and composition of the CEI are affected by the charging and discharging voltages. High voltage will lead to more electrolyte decomposition, causing most of the products to accumulate in the form of a qSEL, while the SEL is relatively thin. Meanwhile, the gas release at high voltage will cause the CEI to become more porous, making it easier for the electrolyte to penetrate. Excessive thickening of the qSEL elevates interfacial impedance, whereas a thin and porous SEL accelerates electrolyte decomposition and drives further qSEL build-up, ultimately precipitating pronounced capacity fade in Ni-rich cathodes under high-voltage operation (Fig. 3d).¹²³

Although interfacial side reactions are usually viewed as detrimental, the literature offers two contrasting perspectives on the CEI. One academic thought posits that the CEI scavenges active Li⁺ and TM species, lowering coulombic efficiency and raising cell impedance.^126–128 Another view is that due to the ionic conduction and electronic insulation properties of the CEI, it can ensure the rapid migration of Li⁺, prevent the continued decomposition of electrolyte components, suppress the co-embedding of solvent molecules and thereby aviods the irreversibly damaging of the electrode material. That is, the CEI can actually play a protective role against further electrolyte erosion and effectively inhibit the continuation of interfacial side reactions.¹²⁹

2.1.3. Surface phase transition. The surface irreversible phase transition is a resultant manifestation of Li/Ni mixing and oxygen release. As Li⁺ continues to detach, the Ni-rich cathode surface undergoes a spontaneous phase transition to an electrochemically inactive rock salt phase due to the lower formation energy of the rock salt and spinel phase transitions compared to the layered structure. During the charging process, Li⁺ is continuously detached and Ni³⁺ is oxidized to the high valence Ni⁴⁺. In order to maintain the charge balance, Ni⁴⁺ would spontaneously reduce to Ni²⁺, with the emission of lattice oxygen and the eventual formation of the NiO phase.¹³⁰

This phase transition is related to the state of charge (SOC) and charge/discharge rate, and intensifies as the SOC and charge rate increase.¹³¹ At low SOC, this irreversible phase transition is achieved by Ni disproportionation and oxygen release (eqn (1) and (2)); at high SOC, especially in the H3 phase, this irreversible phase transition is mainly caused by oxygen release (eqn (3) and (4)).¹³² In addition, this irreversible phase transition occurs during thermal decomposition, intensifies with increasing Ni content, and leads to the release of O₂ and CO₂.^133,134 In the practical application of Ni-rich cathodes, due to the slow solid phase diffusion kinetics of Li⁺, the surface structure of the material will be more likely to form a lithium-poor region and accelerate the structural decline. The phase transitions coupled with lattice-oxygen loss on the surface trigger electrolyte oxidation and vigorous exothermic reactions, severely undermining both cycling stability and safety.^135,136

3Li_xNiO₂ (layered) → Li_xNi₂O₄ (spinel) + NiO (rocksalt) + xLi₂O (1)

3NiO₂ (layered) → Ni₃O₄ (spinel) + 2[O] (3)

Ni₃O₄ (spinel) → 3NiO (rocksalt) + 2[O] (4)

2.1.4. Cross-effects caused by different anodes. Beyond focusing on the intrinsic degradation mechanisms of Ni-rich cathodes, from a full cell perspective, different anode materials can also indirectly influence cathode degradation mechanisms and stability through cross-effects. Specifically, anode byproducts dissolve and migrate to the cathode, then deposit on the surface of cathode, triggering interfacial deterioration of Ni-rich cathodes.¹³⁷ Specifically, when the full cell is composed of graphite, prelithiated SiO_x and lithium metal anode respectively with LiNi_0.8Co_0.1Mn_0.1O₂ (NCM811) cathode, due to the gradually increasing surface reactivity of graphite, SiO_x and lithium metal anode, and the significant differences in electrode expansion and solid electrolyte interphase (SEI) stability, their influence on the degradation of the Ni-rich cathode is also different (Fig. 4a).¹³⁸ Highly reactive anodes will intensify interfacial side reactions, thereby leading to the deterioration of Ni-rich cathode performance. The cathode paired with graphite exhibits the lowest charge transfer impedance (R_ct) as shown in Fig. 4b. As the anode reaction activity increases, the separator after cycling shows a distinct yellowish-brown color, the pores are severely blocked by decomposition products, and the surface of the cathode is also covered by a thick coating. In addition, the higher the anode activity, the thicker the CEI layer on the surface of the Ni-rich cathode after cycling (Fig. 4c), and the interface side reactions will continue to occur inside the CEI (Fig. 4d and e). These phenomena directly demonstrate that an anode with stronger reactivity will intensify the decomposition of the electrolyte, and its products will migrate across the separator and eventually deposit on the surface of the cathode, forming a coating that hinders ion transport.

Fig. 4 (a) Schematic of the temperature-accelerated degradation pathway. Heat amplifies reactivity at both electrodes, leading to increased SEI generation and soluble byproduct formation at the anode. These species then crossover to the cathode, driving excessive CEI formation, with the severity scaling with anode reactivity (graphite, SiO_x, Li-metal). (b) Fitted R_ct values plotted versus specific capacity (curve colors correspond to the paired anode: orange for Gr, purple for SiO_x, blue for Li metal, and green for pristine). (c) Thickness of the CEI layer for cathodes cycled against graphite, SiO_x, and Li metal anodes. (d) Integrated yield comparison of selected secondary fragments. (e) 3D reconstruction of spatial distribution for selected CEI components in cycled cathodes.¹³⁸ Copyright 2025, John Wiley and Sons.

2.2. Damage throughout the structure

2.2.1. Microcracks and granular fracture. Microcracking and particle cracking are caused by interfacial mechanical stresses originating from irreversible phase transitions and oxygen loss.^139–141 Differentiated by location, there are two general types of microcracks: intergranular cracks and intragranular cracks (Fig. 5a).¹⁴² Intergranular cracks are those that occur at the grain boundaries of secondary particles, while intragranular cracks occur within primary particles.

Fig. 5 (a) Schematic diagram of the two cracks.¹⁴² Copyright 2023, Elsevier. (b) Comparison of the c-axis lattice parameters, a-axis lattice parameters, and unit cell volumes for LiNiO₂. For comparison, the converted lattice parameters of the monoclinic phase are plotted together with the hexagonal unit cell parameters.¹⁵⁶ Copyright 2019, Royal Society of Chemistry. (c) Secondary particles consist of (top) randomly oriented, (middle) radially oriented, and (bottom) size-refined primary particles. Depending on the microstructure, the build-up of local stress concentrations and stress distribution is different during charging.¹⁵⁷ Copyright 2019, John Wiley and Sons. (d) 2D mapping of mesoscale SOC heterogeneity in cycled NCM secondary particles.¹⁶⁰ Copyright 2019, John Wiley and Sons.

2.2.1.1. Intergranular cracks. Intergranular cracking is more detrimental to the material than intragranular cracking; the latter originates mainly from the anisotropic shrinkage and expansion of primary particles during irreversible phase transitions.^143–149 The intergranular cracks generated along the grain boundaries hinder the contact of the primary particles and therefore have an adverse effect on the internal conductivity of the secondary particles, which will result in an uneven charging state within the Ni-rich cathode particles.¹⁵⁰ In addition, the generation of intergranular cracks allows the electrolyte to penetrate further into the interior of the particles, where more surface side reactions occur, accelerating the collapse of the Ni-rich cathode structure and the decay of the performance.¹⁵¹ Specifically, there are two factors that contribute to the formation of intergranular cracks: first, the irreversible H2–H3 phase transition.^152–155 This leads to anisotropic lattice strains and sudden changes in lattice parameters within the Ni-rich cathode. As shown in Fig. 5b, the a-axis length and c-axis length of the Ni-rich cathode keep changing with continuous delithiation, and the change in the lattice volume follows the trend of the c-axis length. For the LiNiO₂ cathode, unlike the overall decreasing trend of the a-axis, the c-axis will undergo a sudden change during the H2–H3 phase transition, and this sudden change will lead to localized strain concentration, which will initiate microcracks. In addition, due to the irregular shape of the primary particles of the Ni-rich cathode, their anisotropic deformation due to the phase transition is random.¹⁵⁶ Meanwhile, the irregular arrangement of the primary particles leads to the interaction of forces generated by the expansion and contraction of two neighboring primary particles at their grain boundaries (Fig. 5c). This anisotropic stress and strain can be mitigated by refining the primary particle morphology.¹⁵⁷ The second reason for the formation of intergranular cracks is the different SOC between different primary particles and even within individual active particles.¹⁵⁸ According to previous studies, cracks expand from the core to the surface,¹⁵⁹ which is due to the high degree of reduction on the surface of the Ni-rich cathode, while the interior is highly oxidized and inhomogeneous (Fig. 5d).¹⁶⁰ Firstly, when the SOC is inhomogeneous, a charge compensation effect occurs in the Li-poor regions of the particles, causing Ni²⁺ near Li-vacancies to be oxidized to Ni⁴⁺.^161,162 Similarly, O will also have a charge compensation effect,¹⁶³ forming O vacancies and pores, which leads to microcracks. This also lowers the migration barrier of Ni, leading to easier migration of Ni²⁺ to the Li layer and the formation of disordered spinel and rock salt phases. During the charging and discharging process, neighboring primary particles with layered and disordered phases produce different volume changes, leading to forces at the grain boundaries of the two primary particles, which in turn produce intergranular cracks along the boundaries.¹⁶⁴ In addition, under high SOC conditions, the de-embedding of Li⁺ effectively reduces the shielding effect between neighboring oxygen planes, leading to an increase in the electrostatic repulsion between the oxygen planes, which makes it easier for oxygen to lose electrons and produce O₂ gas with the formation of oxygen vacancies.¹⁶⁵ This is also the reason for the sudden contraction of the c-axis.^70,144,166 Porosity in Ni-rich cathodes increases when O is continuously oxidized and oxygen is produced. Porosity is associated with the nucleation of microcracks and is negatively correlated with SOC, which accelerates the decay of the structure.¹⁵⁹ In addition, there is a rock salt phase region formed around the pores,¹⁶⁷ and the pores can disrupt the diffusion pathway of Li⁺ during cycling, leading to inhomogeneous electrochemical activity of the particles and increasing the likelihood of intergranular cracks.¹⁶⁸ In addition, there is a significant effect of the difference in SOC on the formation of intergranular cracks. At high SOC, the NCM secondary particles tend to rupture and produce a high density of cracks during the initial cycling. The cracks nucleate at the end of the charging process and grow/expand during the discharging process. At low SOC, the NCM secondary particles, on the other hand, exhibit fatigue cracking behavior with chronic degradation.¹⁶⁹
2.2.1.2. Intragranular cracks. Intragranular cracks are another type of crack in Ni-rich cathode materials. The causes of intragranular cracks are relatively complex. The root cause of these cracks is localized stress. As shown in Fig. 6a,¹⁷⁰ specifically, they may be due to fatigue cracking caused by the expansion and contraction of primary particles during charging and discharging. They may also be induced by inhomogeneous expansion and contraction within the crystal as the SOC varies due to TM inhomogeneity. They may also occur due to tensile stresses in the fragile rock salt phase region formed by Li/Ni mixing. In addition, since the diffusion ability of Li⁺ in the rock salt phase generated by interface side reactions is different from that in the layered structure, Li⁺ delithiation is uneven in the extremely lithium deficient region, resulting in a stress field inside the particles. Intragranular cracks occur when the stress exceeds the fracture strength of the crystal plane.

Fig. 6 Schematic representation of the evolution mechanisms of intragranular cracks. (a) Charged NCM single crystal viewed along the (003) plane showing initial intragranular cracking and an initial rock salt-like layer (RSL) on the outer surface (orange). (b) Discharged NCM crystals showing the detrimental effects of intragranular cracking after aging. (c) and (d) Effects of crack reaching the outer surface and flooding with the electrolyte. This leads to the formation of an additional RSL (dark red) and eventually detachment of a segment of the original crystal. (e) and (f) Effects of plane gliding, which also leads to detachment of sections of the original crystal and formation of an additional RSL.¹⁷⁰ Copyright 2024, American Chemical Society. (g) High angle angular dark field (HAADF) images from cycled LiNi_1/3Mn_1/3Co_1/3O₂ cathode particles, showing intragranular cracks along the (001) plane. The yellow arrows indicate real cracks and the pink arrows indicate incubation cracks. Scale bars: 10 nm. (h) Transmission electron microscope (TEM) image of a damaged primary particle in the cycled NCM cathode with diffraction patterns (inset) from the marked regions.¹⁶⁴ Copyright 2016, John Wiley and Sons. (i) HAADF images overlaid diagram showing the apparent dependence of intragranular cracking on the cycle voltage; when cycled below 4.5 V, intragranular cracks can be hardly generated, while above 4.7 V, the density of intragranular cracks shows a drastic increase. (j) Schematic diagrams illustrating the dislocation-assisted crack incubation, propagation and multiplication process. Intragranular cracking as a critical barrier for high-voltage usage of layer-structured cathodes for LIBs.¹⁴⁷ Copyright 2017, Springer Nature. (k) Scanning diffraction X-ray microscopy (SDXM) images of the (003) peak at OCV, charged to 4.1 V, and charged to 4.5 V. The d-spacing variation is expressed by blue and red colors. The length and orientation of the quiver represent the magnitude and direction of lattice rotation, respectively. A fixed scale is used for all three figures. (l) Lattice rotation magnitude. (m) Schematic illustration of lattice rotation.¹⁷³ Copyright 2024, The American Association for the Advancement of Science. (n) and (o) Scanning transmission electron microscopy (STEM) images of single-crystalline NCM at the 4.4 V charge status (cycled in a full cell between 2.7 and 4.4 V for 120 cycles). (p) and (q) STEM images of single-crystalline NCM at the discharge status (cycled in a full cell between 2.7 and 4.4 V for 120 cycles). The red arrows indicate the gliding marks inside the single crystal. (r) COMSOL-simulated shear stress along the yz direction during charge (delithiation). (s) COMSOL-simulated shear stress along the yz direction during discharge (lithiation). (t) Schematic illustration of the structural evolution of single-crystalline NCM upon cycling.⁶⁵ Copyright 2020, The American Association for the Advancement of Science.

Furthermore, at the atomic scale, Lin et al.¹⁴⁶ found that Li/Ni anti-site defect-rich regions with lattice distortions are the nucleation sites for intragranular cracks in primary particles. The strain difference between the layered phase and the rock salt phase and the coulomb repulsion in the cation-rich region are the two driving forces for the growth of intragranular cracks, which originate from the Li/Ni anti-site defects in the transformed structure. Moreover, in terms of specific elemental considerations, Guo et al.¹⁷¹ found that Co is responsible for the intensification of the intragranular cracking phenomenon. In the deeply charged state, Co exacerbates the transition from the layered phase to the rock salt phase and the accumulation of non-homogeneous strain, which ultimately drives intragranular cracking in the rock salt phase domain along the (003) plane at the Ni-rich cathode. In addition, oxygen vacancy formation and their subsequent migration along the (003) plane introduce pronounced lattice strain and expansion, which in turn promote the propagation of intragranular cracks. Unlike intergranular cracks, intragranular cracks begin inside the grain and the size of intragranular cracks is smaller than that of intergranular cracks, but the number of cracks is larger, so they are able to expose more of the fresh material surface in contact with the electrolyte. In addition, it is worth noting that intragranular cracking is not only a mechanical failure but also a structural degradation.¹⁴⁷ According to the crack generation mechanism and morphology, intragranular cracks can be classified into three types: (1) classic cracks (indicated by yellow arrows in Fig. 6g);¹⁴⁷ (2) dark contrast stripes (indicated by pink arrows in Fig. 6g); and (3) fine-serried low angle grain boundary (white circles in Fig. 6h).¹⁶⁴

Currently, the first two types of intragranular cracks are believed to be formed due to heterogeneous delithiation. The internal structural properties of dark contrast stripes were analyzed by Yan et al. It was confirmed that the dark contrast stripes are premature classic cracks, and the authors constructed a model of the crack formation process.¹⁴⁷ According to the model (Fig. 6i and j), in the first stage, the damage to the structure by cracks becomes progressively more severe with the increase of voltage because the creation and development of cracks are electrochemically driven. In the second stage, dislocations act as nucleation sites for intragranular crack hatching, initiating premature cracks. In the third stage, premature cracks gradually develop into real cracks with two parallel planes (classic cracks). The third type of intragranular cracks (fine-serried low angle grain boundary) is mainly formed by mechanical damage caused by the concentration of stress within the primary particles. The main causes of this mechanical degradation are (1) fatigue strain during cycling, (2) dislocations between primary particles and (3) “false” primary particles with grain boundaries.¹⁷²

Additionally, inhomogeneous delithiation induces lattice strain alongside lattice rotation (refers to displacement of the momentum transfer in the 3D reciprocal space in directions orthogonal to the displacement caused by lattice strain). Moreover, the degree of lattice rotation intensifies as the voltage increases as shown in Fig. 6k–m, ultimately leading to mechanical degradation of Ni-rich cathodes in the form of partial irreversible plane gliding and intragranular cracking.¹⁷³ Compared to polycrystalline particles, lithium diffusion paths are longer in larger single-crystal cathodes,¹⁷⁴ making inhomogeneous delithiation more severe. Consequently, lattice rotation is more pronounced in single-crystal cathodes. This lattice rotation cannot be eliminated by lithium reinsertion like lattice strain, and it is largely unavoidable since it serves as one of the fundamental mechanisms for distorting the lattice to accommodate heterogeneous electrochemical reactions within single-crystal cathodes during lithium extraction–insertion cycles. Although lattice gliding along the (003) crystal plane is largely reversible, lattice gliding and microcracks were observed on the surface of the Ni-rich cathode in the charged state. These microcracks originated from within the crystal. Most intracrystalline cracks and lattice gliding disappeared in the fully discharged state, with only minor traces of the gliding process remaining (Fig. 6n–q). However, non-self-recovering lattice gliding and microcracks persist, as confirmed by COMSOL simulation results (Fig. 6r and s). The shear stress component along the yz direction can trigger gliding along the (003) plane. During charging/discharging, the sign of shear stress reverses while its absolute value is not equal, leading to incomplete reversibility of lattice gliding. This partially irreversible slip accumulates into crack openings after prolonged cycling, consequently causing post-cycling surface cracking in single crystals (Fig. 6t).⁶⁵

Based on the formation and evolution mechanism of intragranular cracks, Morzy et al.¹⁷⁰ concluded that intragranular cracks negatively affect the battery performance in two main ways, as shown in Fig. 6b–f. First, for intragranular cracks that are not completely sealed (i.e., fail to self-heal) during discharge and re-formed during subsequent charging, they will expand and reach the surface of the primary crystal after multiple cycles, leading to electrolyte flooding, surface reduction (Fig. 6c) and material detachment (Fig. 6d). Secondly, the repeated formation and closure of intragranular cracks can lead to (003) plane gliding (Fig. 6e), which may even lead to the fragmentation of primary grains and the formation of flakes after long-term cycling (Fig. 6f).

2.2.1.3. Detrimental effects of microcracks. The generation of microcracks is the main reason for the decrease in the cycling performance of Ni-rich and low cobalt cathodes. Intergranular cracks will make the particles fail to connect and thus reduce the conductivity of secondary particles,¹⁵⁰ and also exacerbate the uneven SOC, which further reduces the activity of the material and makes the particles more susceptible to chalking and fragmentation. In addition, the formation of cracks will expose more of the inner surface of the active material, accelerate the penetration of the electrolyte into the interior of the particles and result in the dissolution of TMs and the loss of active substances.^148,175 At the same time, the electrolyte reacts with unstable Ni⁴⁺ in the material to form a disordered phase.¹⁷⁶ Due to the significant differences in the crystal lattice between the surface rock salt phase, the spinel-like phase, and the layered phase in the bulk phase, the cracking of secondary particles is further exacerbated.^66,177 Ultimately, it will lead to hindered diffusion of Li⁺, increased impedance, and decreased electrochemical performance. This side reaction causes phase transition of the layered cathode, leading to structural collapse and O₂ release.^178–180

2.2.2. Cation mixing. Cation mixing, also known as Li/TM ion mixing, in Ni-rich cathodes mainly refers to the phenomenon where transition metal ions Ni²⁺ and Li⁺ occupy each other's position in Ni-rich cathode materials. The cation mixing phenomenon exists not only in the synthesis process of battery materials, but also throughout the battery life.
2.2.2.1. Causes of cation mixing. The reasons for the occurrence of Li/Ni mixing in Ni-rich cathodes are complex. The main reasons include (1) similar ionic radius, (2) low migration energy barriers of Ni ions, (3) magnetic interactions, and (4) thermodynamic instability. Generally, layered materials have the R3̄m structure as shown in Fig. 7a.¹⁸¹  In this structure, oxygen anions form a cubic close-packed lattice, with lithium and transition metal (TM) cations orderly occupying alternating layers of octahedral sites along the crystallographic c-axis ([001] direction).¹⁸² Therefore, the TM sites (3a) and Li sites (3b) of the standard R3̄m structure are clearly separated. However, during the sintering process of Ni-rich cathode materials, Ni²⁺ is difficult to be completely oxidized to Ni³⁺. Given the similar radius of Ni²⁺ (0.069 nm) and Li⁺ (0.076 nm), during high-temperature synthesis, the spontaneous occupation of Li sites by Ni²⁺ significantly increases the system's configurational entropy (ΔS > 0). Driven by high temperature (T), the entropy increase (−TΔS) dominates the system's Gibbs free energy (ΔG), resulting in ΔG < 0. Therefore, the formation of Li/Ni antisite defects constitutes a thermodynamically spontaneous process, primarily driven by the thermodynamic principle of entropy increase. This ultimately leads to Ni deviations from stoichiometric ratios in Ni-rich cathodes (as exemplified by LiNiO₂, where the chemical formula changes from ideal LiNiO₂ to non-stoichiometric Li_1−xNi_1+xO₂).^183,184 Therefore, excess lithium is often required during synthesis to counteract this thermodynamic tendency. The excess lithium increases the lithium chemical potential and, following the law of mass action, preferentially occupies lithium layer sites.¹⁸⁵ As a result, Ni²⁺ is confined to the transition metal layer, effectively reducing cation mixing and ensuring the integrity of the layered structure.

Fig. 7 (a) The illustration of the well-ordered R3̄m structure. (b) The illustration of the R3̄m structure with Li vacancies in the highly charged state. (c) The illustration of the partially cation mixed phase with TM ions in the Li slab. (d) The illustration of the cation disorder or cation mixing phase with Fm3̄m structure.¹⁸¹ Copyright 2015, John Wiley and Sons. (e) and (f) Migration routes of Ni ions to Li sites in the layered structure: (e) O_h → T_d → O_h route and (f) O_h → V_O → O_h route. (g) and (h) Energy barriers calculated for migration of a Ni ion for (g) the O_h → T_d → O_h route and (h) the O_h → V_O → O_h route.⁶⁹ Copyright 2011, American Chemical Society. (i) and (j) Nudged elastic band (NEB) calculations comparing the energy barriers for Ni migration in LiNiO₂ at (i) 25% delithiation under pristine and oxygen-deficient conditions and (j) 50% delithiation (green, grey, and red spheres represent Ni, Li, and O ions, respectively).¹⁸³ Copyright 2025, Elsevier. (k) Influence of oxygen defect on structural instability. (l) Thermogravimetric analysis demonstrating oxygen deficiency at low oxygen partial pressure.¹⁸⁷ Copyright 2015, Elsevier. (m) Direct proportionality between oxygen defect ratio and cation mixing degree.¹⁸⁸ Copyright 2016, Elsevier. (n) and (o) Two inequivalent positions in the Li layer: (n) configuration A and (o) configuration B. The grey spheres denote Ni, the blue spheres denote Co, the green spheres denote Li and the red spheres denote O.¹⁹⁰ Copyright 2014, Royal Society of Chemistry.

In addition, a drive for Li/Ni mixing is also generated during Ni-rich cathode cycling, which is attributed to the low migration barrier for Ni ions. In the charging process, as Li⁺ is released (Fig. 7b), the migration barrier for Ni ions is reduced and the TM ions occupy the Li sites (Fig. 7c). The material would gradually become a disordered spinel phase and collapse during charge and discharge, and eventually transition from the original layered phase to the rock salt phase (Fig. 7d).¹⁸¹ There are two migration paths for Ni during cycling.^69,186 Path a (Fig. 7e): Ni ions migrate from one octahedral (O_h) site to another octahedral site (O_h–T_d–O_h) via a tetrahedral (T_d) site. This migration may occur within or between layers. Path b (Fig. 7f): Ni ions migrate from one octahedral (O_h) site to another octahedral site (O_h–V_O–O_h) via vacancies (oxygen vacancy, V_O). This migration occurs within the TM layer. Based on the calculations of the migration barriers for the two pathways, it is found that the energy barriers in the O_h–T_d–O_h route are related to the cation configuration (Fig. 7g). In addition, by comparing the number of nearest cations and their distances from the T_d and O_h (Li) sites, as well as the magnitude of the energy barriers at these positions, Kim et al.⁶⁹ confirmed that the main factor determining the migration energy barriers is the coulomb repulsion from the surrounding cations. On the other hand, for the O_h–V_O–O_h pathway, the migration energy barrier is very little affected by the configuration (Fig. 7h). In conclusion, the migration probability of both routes is the same in the Li 100% and Li 50% (configuration 1 containing more Li⁺) states, whereas Ni ions are more prone to migrate via the O_h (Ni)–T_d–O_h (Li) route in the Li 50% (configuration 2 containing less Li⁺) and Li 0% states, which may lead to phase transition and performance degradation. In addition, according to the stabilization energy of octahedral sites, Ni ions tend to migrate from octahedral sites to tetrahedral sites. Li et al. further confirmed that the Ni migration barrier is influenced by Li vacancy concentration and oxygen vacancy concentration through density functional theory (DFT) calculations. When 25% and 50% of Li ions were extracted from a Li₁₂Ni₁₂O₂₄ supercell without oxygen vacancies, the energy barrier for Ni ion migration to Li layer octahedral sites decreased from 3.20 eV to 2.74 eV (Fig. 7i and j) as Li vacancies formed, indicating that the presence of Li vacancies promotes Ni migration. Furthermore, the calculations indicate that oxygen vacancies exert a more pronounced effect on Ni migration in Ni-rich cathodes. The presence of oxygen vacancies reduces the Ni migration barrier (2.38 eV) and formation energy (0.61 eV), suggesting that oxygen vacancies also promote Ni migration.¹⁸³ This result is also supported by Bi et al.,¹⁸⁷ who found that oxygen vacancies also form in Ni-rich cathodes when the partial pressure of oxygen during sintering is too low, thereby promoting the occurrence of Li/Ni mixing (Fig. 7k–m).¹⁸⁸ Based on the above findings, the researchers further found the relationship between cation mixing and oxygen release. With the release of oxygen during charging, more oxygen vacancies are generated in the Ni-rich cathode, which promote the migration of Ni ions via O_h–V_O–O_h and continuously intensify Li/Ni mixing during long cycling.¹⁸⁹

The arrangement of cations in the lattice is also related to the nature of the magnetic properties. Fig. 7n and o show two configurations of the atomic arrangement in the lattice: configuration A: six pairs of 180° Ni–O–Ni structures, and configuration B: six 90° Ni–O–Ni structures.¹⁹⁰ According to the super-exchange interaction,¹⁹¹ Ni²⁺ (t⁶_2ge²_g) and Ni³⁺ (t⁶_2ge¹_g) have fully filled t_2g states and partially filled e_g states, respectively. Therefore, the 180° Ni–O–Ni super-exchange interaction is stronger than the 90° Ni–O–Ni super-exchange interaction. That is, when Ni–O–Ni is in the 180° configuration, more energy is gained through orbital interactions than when it is in the 90° configuration. Therefore, 180° Ni–O–Ni can stabilize Ni in the Li layer. Further studies by Zheng et al.¹⁹² showed that Ni²⁺–O–Ni²⁺ has the strongest super-exchange interaction, and that the more Ni²⁺–O–Ni²⁺ there is, the lower the formation energy of Li/Ni mixed defects, and the easier Li/Ni mixed defects can be formed. In general, when there is no Li/Ni mixing, the Ni in the TM layer and the Ni in the neighboring layer form a 90° Ni–O–Ni structure. When Li/Ni mixing occurs, the randomly distributed Ni in the Li layer may form an 180° Ni–O–Ni super-exchange structure.

In addition, it is shown by studies that the occurrence of cation mixing in Ni-rich materials is essentially due to their poor thermal stability. Taking LiNiO₂ as an example, its decomposition occurs during the heating process, which ultimately leads to Li/Ni mixing and the formation of the NiO rock salt phase. The thermal decomposition equations are eqn (1) and (2).

2.2.2.2. Effects of cation mixing. First, cation mixing would directly lead to lattice oxygen release. Li et al.¹⁸³ confirmed through DFT calculations that during Li/Ni cation mixing, Ni ions migrate into the Li layer, leaving vacancies in the TM layer. These TM vacancies may be filled by Li⁺ from the Li layer to form a Li–O–Li structure, or may remain vacant to form a Li–O–Vac_TM structure. Both structures promote oxygen oxidation reactions.^193,194 Furthermore, for Ni-rich cathode systems with different Li/Ni mixing ratios, the Bader charge, magnetic moment, and oxygen vacancy formation energy were calculated at different charge states (Fig. 8a–c). Fig. 8a and b reveal that oxygen atoms in the Li–O–Vac_TM configuration exhibit smaller overall Bader charges and higher magnetic moments, indicating a greater tendency to participate in charge compensation during lithium extraction and a higher oxidation state. Furthermore, regardless of the configuration, the oxygen vacancy formation energy in Ni-rich cathodes with Li/Ni mixing is lower than in the perfect layered structure (0% Li/Ni mixing; Fig. 8c). Ab initio molecular dynamics (AIMD) simulations reveal that in systems with Li/Ni mixing, oxygen atoms adjacent to TM vacancies form O–O dimers at lower delithiation concentrations, whereas this does not occur in the perfect layered structure. This indicates that cation mixing accelerates the extraction of lattice oxygen. In addition, refinement results from neutron diffraction provide further support for these simulation findings. As shown in Fig. 8k,¹⁹⁵ during delithiation, the degree of Li/Ni mixing correlates with trends in the isotropic temperature factor of oxygen atoms (representing lattice oxygen instability) and lattice stress (leading to structural collapse). This indicates that the activation of anion redox reactions and the abrupt increase in stress are closely related to Li/Ni mixing, ultimately leading to structural degradation and reduced cycling life in high-voltage Ni-rich cathodes.

Fig. 8 (a) Bader charge and (b) oxygen magnetic moment evolution with delithiation. (c) Li/Ni mixing dramatically lowers the oxygen vacancy formation energy, confirming its role in promoting oxygen release.¹⁸³ Copyright 2025, Elsevier. (d) Schematics showing that cation-mixed Ni²⁺ contributes to structural integrity retention during discharge and charge, highlighting its role in stabilizing the crystal framework.¹⁹⁷ Copyright 2017, The Royal Society of Chemistry. Schematic representation of (e) layered-ordered Li_xNCM and (f) the local Li-rich and TM-rich domain created after Li/Ni disordering associated with the potential spontaneous electronic reshuffling. (g) Spin density diagram of the local disordered domain, in which the Ni/Li disorder is in the form of Ni_Li–Li_Ni antisites with x = 1, 0.67 and 0.33, and Ni_V–V_Ni antisites with x = 0. The yellow and cyan isosurfaces represent the electrons with spin up (↑) and down (↓), respectively. The black arrow reflects the electronic reshuffling induced by cation disorder compared to that in the layered-ordered system. The formal charge reaction equations during disordering are marked under the diagrams. (h) Formation energy of V_O after cation disorder as well as (i) the formation energy of cation disorder after oxygen loss. (j) Schematic illustration of the self-sustaining degradation process. (k) Evolution of Ni migration extent, oxygen isotropic displacement parameter (U_iso), and structural stress with cycling, as refined from operando neutron diffraction data. (l) Schematic of the dual impact of antisite defects on structural evolution.¹⁹⁵ Copyright 2024, John Wiley and Sons.

Chen et al.¹⁹⁶ elucidated the mechanism of lattice oxygen release induced by cation mixing at the electronic scale under highly delithiated conditions. This mixing process induces electronic structure reorganization in the high delithiation state as shown in eqn (5) (Fig. 8e–g). The TM ions migrate to the Li layer and are reduced and leave behind electronic holes, while the electronic states near the Fermi energy level will then get holes. During deep delithiation, this electronic recombination excites lattice oxygen activity and increases the risk of lattice oxygen loss. Further thermodynamic calculations show that regardless of the coordination mode (O(A): Li–O–Li; O(B): Li–O–TM; O(C): TM–O–TM), there is a mutually inverse thermodynamic relationship between lattice oxygen stability and cation mixing (Fig. 8h and i). Cation mixing induces lattice oxygen instability and promotes oxygen loss and oxygen vacancy formation, a process that in turn exacerbates cation mixing and further electronic recombination. This “self-sustained” degradation pathway continues throughout the structural decline of the Ni-rich cathode, with cation mixing as the nucleation site and cation mixing and oxygen loss processes occurring alternately. O(A)/O⁽²⁻ⁿ⁾⁻ ions in the lattice accumulate and eventually escape from the lattice, while residual O(C) ions and migrating TM ions accumulate to form the rock salt phase, which ultimately leads to an extensive irreversible phase transition (Fig. 8j). In other words, cation mixing is closely related to lattice oxygen stability.

Ni(3a)^x+ + (x − 2)e⁻ → Ni(3b)²⁺, nA^y+ + (x − 2)h˙ → nA^(y+x−2)/n+ (5)

To further discuss the impact of Li/Ni mixing on Ni-rich cathodes, it is also necessary to understand how Li⁺ migrates. There are two main Li⁺ migration paths:¹⁹⁸ the first type of migration path (Fig. 9a) occurs when the two Li sites immediately adjacent to the endpoints of the hop are simultaneously occupied by Li⁺. The diffusing Li⁺ then migrates along the shortest path connecting the initial site of the hop and the vacancy. This path passes through a dumbbell of the oxygen ion, called the oxygen dumbbell hop (ODH). The second type of migration path (Fig. 9b), called the tetrahedral site hop (TSH), is a situation where Li⁺ migrates along a curved path through a tetrahedral site when one or both sites immediately adjacent to the hop endpoints are vacant. Therefore, under the influence of spatial effects and electrostatic interactions, when Li/Ni mixing occurs, the activation energy of Li⁺ migration through the TSH pathway increases, and the spacing of Li layers decreases, which ultimately inhibits the migration of Li⁺.¹⁹⁹ Zhao et al.²⁰⁰ have shown that the Li layer spacing decreases (2.66 Å, 2.65 Å, 2.63 Å) with the increasing Li/Ni mixing degree (0%, 3.7%, 7.4%). The calculated Li⁺ diffusion barriers were 0.496 eV and 0.608 eV for the models with 3.7% and 7.4% Li/Ni mixing degree, respectively. The migration coefficients of Li⁺ decreased with the increase of Li/Ni mixing (Fig. 9c). In addition, the migration barrier of Ni ions after the occurrence of Li/Ni mixing and dislocation was as high as 1.05 eV, indicating that it is difficult to migrate (Fig. 9d). Ni²⁺, located in the Li layer, has a higher valence state than Li⁺ and therefore generates greater electrostatic repulsion. As a result, Li⁺ migration along both the ODH and TSH paths will be inhibited, which means that Li/Ni mixing inhibits the migration of Li⁺.

Fig. 9 The two lithium migration paths in layered Li_xCoO₂. The filled circles are lithium ions, the empty squares are lithium vacancies, the large empty circles are oxygen ions, and the small empty circles are cobalt ions. (a) The ODH occurs when the sites a and b adjacent to the end points of the hop are simultaneously occupied by lithium ions. (b) The TSH occurs when one or both of the sites adjacent to the endpoints of the hop are vacant.¹⁹⁸ Copyright 2001, Elsevier. (c) The energy curves of Li atom migration paths for different Li⁺/Ni²⁺ ion exchange rates. (d) The energy curve of the Ni atom migration path. The insets are illustrations of the migration path.²⁰⁰ Copyright 2017, Royal Society of Chemistry. (e) Surface morphology of NCM after 30 h of electrolyte exposure (equivalent to one full cycle duration in this study). (f) Surface structure of NCM after 1 cycle (2.0–4.7 V), with the blue arrow indicating the surface reconstruction layer. (g) Fast Fourier transform (FFT) pattern from the surface reconstruction layer (Fm3̄m [110] zone axis). (h) FFT pattern from the NCM bulk structure (R3̄m [100] zone axis). (i) Orientation-dependent thickness variation of the surface reconstruction layer of NCM after 1 cycle (2.0–4.7 V). (j) Atomic-scale image showing loose surface layers of NCM, with blue lines demarcating the interface between bulk NCM and the reconstruction layer. The blue lines indicate the boundaries between the NCM layered structure and the surface reconstruction layer in all images. Scale bars: 2 nm in (e)–(j).²⁰¹ Copyright 2014, Springer Nature. (k) HAADF STEM images at different regions from the bulk to the surface of Ni-rich cathodes and the corresponding structure models for explaining the structure transition process. The white dashed lines denote the surface reconstruction layer at different stages.²⁰² Copyright 2018, Elsevier.

Lin et al.²⁰¹ found that Li/Ni mixing was evident on the surface of the Ni-rich cathode (Fig. 9e) and thickened after cycling (Fig. 9f–h). Rock salt phases were formed in the diffusion path of Li⁺ (Fig. 9i), which would seriously hinder the migration of Li⁺. In addition, a loose atomic layer exists on the outer surface of the mixing layer (Fig. 9j). Lin et al.²⁰² found that after cycling, the material is layered inside while the surface is in a rock salt phase, and the transition region exhibits a disordered layered phase and a defective rock salt phase. As shown in Fig. 9k, they similarly suggested that Ni migrates along the Li diffusion direction during cycling, which leads to the evolution of the structure and dissolution of Ni, ultimately leading to a decrease in capacity. In summary, Li/Ni mixing is believed to be the origin of the phase transition. This irreversible phase transition also exacerbates oxygen release, promotes interfacial reactions and contributes to the degradation of the electrochemical performance of Ni-rich cathodes.²⁰³

Although most of the studies showed that Li/Ni mixing is harmful to the overall performance of Ni-rich cathodes, some of them found that a moderate amount of Li/Ni mixing is beneficial to Ni-rich cathodes. Tang et al.²⁰⁴ investigated the relationship between Li/Ni mixing and Li⁺ diffusion barriers, and found that when Ni²⁺ ions occupy Li⁺ sites, the stronger interaction between Ni and O drives the nearest Li⁺ ions to diffuse into the neighboring vacancies. In other words, Li/Ni mixing will promote the diffusion of Li⁺ and enhance the rate performance of Ni-rich cathodes. Moderate cation mixing enhances Li⁺ diffusion kinetics and structural stability by facilitating concerted Li⁺ hopping through Ni²⁺-induced vacancy attraction and buffering interlayer repulsion during deep delithiation, thereby improving rate capability and high-voltage cycling retention.^86,205 However, a large amount of Li/Ni mixing could lead to the collapse of the layered structure. Excessive mixing obstructs Li⁺ channels, traps lithium in transition metal layers, and accelerates oxygen loss, triggering irreversible layer-to-spinel/rock-salt phase collapse.²⁰⁶ This is similarly supported by the study conducted by Sun et al.,¹⁹⁷ which found that proper cation mixing reduces repulsive forces and inhibits structural collapse, which is essential for supporting the layered structure of lithium transition metal oxide materials, especially in the deep delithiation state. The above-mentioned gain effect from cation mixing arises because moderate cation mixing enhances structural stability by supporting the Li layer, reducing repulsive forces between adjacent oxygen layers during delithiation, and suppressing transition metal ion migration and phase transitions.²⁰⁷ Ni²⁺ ions occupying Li sites generate a “pillar effect” through electrostatic repulsion,²⁰⁸ effectively preventing the continuous migration of transition metal ions toward lithium sites during charge–discharge cycles (Fig. 8d). In summary, as shown in Fig. 8l, Ni intercalated in the Li layer supports the layered structure but inhibits Li ion diffusion, while Li intercalated in the TM layer forms a Li–O–Li configuration. This configuration triggers anionic (O²⁻) redox activity and induces oxygen release, leading to electrochemical performance degradation and structural deterioration. The concurrently formed oxygen vacancies further promote Ni migration from the TM layer to the Li layer, exacerbating Li/Ni mixing. Notably, the Li–O–Vac_TM configuration can form during Ni migration, which also promotes anion redox activity and disrupts lattice oxygen stability.

3. Gas-generation mechanisms of Ni-rich cathodes

Gas generation in Ni-rich cathodes mainly occurs due to three aspects: (1) the decomposition reactions of surface RLCs, such as LiOH, Li₂CO₃ and LiHCO₃, (2) the release of lattice oxygen at high voltage, and (3) oxidation reactions of the electrolyte with the cathode. The gas-generation mechanism of LIBs is complicated with many influencing factors. The mechanisms and solution strategies should be investigated clearly to reduce the damage caused by gas generation in batteries, such as battery expansion and even explosion.^209,210

3.1. Decomposition of surface residual lithium compounds

Li₂CO₃, which is the main component of RLCs, is not stable, and its decomposition mechanism is mainly divided into chemical decomposition²¹¹ and electrochemical decomposition.^212,213 Freiberg et al.²¹⁴ concluded that the decomposition of surface Li₂CO₃ is a purely chemical decomposition reaction. The decomposition of Li₂CO₃ was demonstrated to be a chemical process by on-line electrochemical mass spectrometry using an electrode consisting only of Li₂CO₃ and conducting carbon. This reaction occurs via protons formed during anodic oxidation of trace alcoholic electrolyte impurities and the electrolyte solvent or induced by high HF in the electrolyte (Fig. 10a), ultimately releasing CO₂ gas. In the field of lithium-air batteries, Mahne et al.²¹³ found that Li₂CO₃ decomposes to form singlet oxygen ¹O₂ at >3.8 V, which is an electrochemical decomposition. By ¹⁸O isotopic labelling of the small amount of surface Li₂CO₃ remaining after synthesis, Renfrew et al.²¹⁵ found that the degradation of the carbonate electrolyte during the first charge process occurred with minimal gas evolution. Most of the CO/CO₂ release during the first charge regardless of the voltage interval is due to Li₂CO₃ oxidation and not because of electrolyte oxidation. In addition, the decomposition of RLCs affects the release of lattice oxygen. When some of the Li₂CO₃ is washed off the surface of the transition metal oxide cathode, CO₂ evolution decreases and lattice-oxygen release is suppressed, resulting in reduced O₂ generation (Fig. 10b). Hatsukade et al.²¹⁶ investigated the origin of CO₂ during cycling of NCM cathodes using the ¹³C isotope labelling method and found that the decomposition of surface Li₂CO₃ at high voltage is one of the reasons for CO₂ generation. However, it is not the main reason. After a long time of high potential cycling, Li₂CO₃ will be consumed continuously, and the CO₂ produced by chemical oxidation and electrochemical oxidation of the electrolyte solvent at high voltage becomes the main reason, and this decomposition is mainly related to the release of lattice oxygen from the NCM (Fig. 10c). Kaufman et al.¹⁰² fabricated a cathode with lithium carbonate as the only active material, and confirmed that the source of CO₂ is the electrochemical oxidation of Li₂CO₃ rather than chemical oxidation by an isotopic labeling method and carbon cathode control experiments. This process is accompanied by the production of reactive oxygen species, which react directly or indirectly with the electrolyte to further produce CO₂ and the rest of the organic debris (Fig. 10d). Regarding the electrochemical decomposition mechanism, Cao et al.²¹⁷ found that Li₂CO₃ would decompose electrochemically at about 4 V when a LiTFSI/LiFSI based electrolyte is used, releasing CO₂ gas and highly reactive monoclinic oxygen ¹O₂. The ¹O₂ would in turn react with the carbon at the cathode and with the organic electrolyte to produce CO₂ and CO gases (Fig. 10e).

Fig. 10 (a) Scheme of Li₂CO₃ decomposition in the LIB environment. Protons catalyze the decomposition of Li₂CO₃, whereas their formation strongly depends on the purity and kind of solvents used.²¹⁴ Copyright 2020, Elsevier. (b) The main mechanisms of the CO/CO₂ release during the first charge.²¹⁵ Copyright 2017, American Chemical Society. (c) The mechanisms of CO₂ production during cycling of NCM cathodes.²¹⁶ Copyright 2018, American Chemical Society. (d) Reactive oxygen released from the NCM lattice is triggered by the oxidation of surface lithium carbonate.¹⁰² Copyright 2021, American Chemical Society. (e) Li₂CO₃ is electro-oxidized to form CO₂ and singlet O₂ (¹O₂). A possible pathway is proposed in the dashed line box. ¹O₂ is highly reactive and it attacks the electrolyte and the carbon substrate to form carbon monoxide and carbon dioxide, which contribute to ∼20% of the overall gas evolution.²¹⁷ Copyright 2022, Springer Nature.

3.2. The release of lattice oxygen at high voltage

Layered oxide cathodes show evident oxygen evolution at high SOC, where increased Ni content correlates with a reduced voltage threshold for lattice oxygen release.^{28,45,218,219} It ultimately triggers structural collapse, parasitic interfacial reactions, and catastrophic thermal runaway. Therefore, understanding the mechanisms that link oxygen activity to gas evolution is pivotal for the rational design of next-generation Ni-rich cathodes.

The mechanism of lattice oxygen release can be explained at two levels: (1) during deep delithiation, overlapping of transition metal (TM) 3d and O 2p orbitals induces anion redox reactions. (2) The unstable layered structure accelerates lattice oxygen release.²²⁰ Among these, the anion redox in the deeply delithiated state constitutes both the thermodynamic driver and the fundamental mechanism for oxygen release from the Ni-rich cathode lattice. Specifically, this refers to the elevation of the Ni⁴⁺/Co⁴⁺ 3d energy levels under highly delithiated conditions, leading to progressive overlap with the O 2p band (as shown in Fig. 11a).²²¹ This p–d orbital overlap enables the rapid absorption of electrons from O²⁻ by the conduction bands of Co and Ni, leading to partial reduction of Ni⁴⁺/Co⁴⁺ and the formation of the oxidized species Oα⁻ (α < 2). This process ultimately drives oxygen out of the lattice, generating reactive oxygen species (O₂ or O⁻) and creating oxygen vacancies.^222,223In situ X-ray absorption fine structure (XAFS) measurements at the Ni K-edge provide direct evidence for this mechanism. As shown in Fig. 11b, during charging, the Ni K-edge of the Ni-rich cathode shifts toward higher energies below 4.2 V, indicating an increase in the Ni valence state. Subsequently, charging to 4.6 V results in a decrease in the Ni valence state due to oxygen release. Furthermore, in the total electron yield mode of the O K-edge XAFS (Fig. 11c), a sharp increase in the E_g state at 4.6 V indicates anion redox participation in charge compensation. With continued lithium extraction, a distinct oxygen redox signature appears near 523.7 eV in the resonant inelastic X-ray scattering (RIXS) spectrum, indicating irreversible oxygen redox reactions within the Ni-rich cathode as shown in Fig. 11d.²²⁴

Fig. 11 (a) Schematic illustration of transition metal levels and oxygen release from the Ni-rich cathode lattice.²²¹ Copyright 2025, Springer Nature. (b) Ni K-edge XANES of charged LiNi_0.95Co_0.05O₂ (NC95). (c) O K-edge XAFS during NC95 charging. (d) O K-edge RIXS of NC95.²²⁴ Copyright 2024, Springer Nature. (e) In situ Raman spectra and (f) the corresponding charge-discharge profile of NCM811 during the third cycle.⁷⁶ Copyright 2020, American Chemical Society. (g) HAADF image showing the EELS scanning path from the particle interior (upper) to the surface (lower). (h) Corresponding O K-edge EELS spectra, revealing the progressive change in the oxygen oxidation state across different regions.²²⁶ Copyright 2022, Elsevier. (i) Schematic illustration of three electrolyte decomposition pathways on the surface of Ni-rich cathodes, including chemical dehydrogenation, chemical oxidation with singlet oxygen, and electrochemical oxidation.²³⁵ Copyright 2023, John Wiley and Sons. (j) Through-transmission ultrasonic scanning image of the formation process of pouch cells with different water content in the electrolyte.²³⁶ Copyright 2022, Chemical Industry Press. (k) Schematic illustration of gassing related processes for an SLP30 graphite/NCM full-cell in a one-compartment configuration and an NCM/Li half-cell in an Al-sealed two-compartment cell. In both cases, electrolyte oxidation generates CO₂/CO and protic decomposition products abbreviated as R-H⁺; however, only in the one-compartment cell, these protic species can diffuse to the anode and get reduced under the release of H₂.²³⁷ Copyright 2016, Institute of Physics.

The instability of the layered structure constitutes the structural cause for lattice oxygen desorption. This instability triggers and exacerbates oxygen loss through multiple mechanisms under high voltage. Firstly, in a deeply delithiated state, highly oxidized TM ions attract electrons from the TM–O bond.²²⁵ Simultaneously, the electrostatic repulsion between TMO₆ layers diminishes, weakening both TM–O and O–TM–O bond strengths. As shown in Fig. 11e and f, during charging, the E_g and A_1g band intensities of NCM811 significantly decay, particularly in the voltage range of 4.15–4.3 V. It indicates that the TM–O and O–TM–O bond strengths weaken at high voltages, making surface oxygen atoms more prone to release from the surface structure.⁷⁶ Furthermore, the crystal structure undergoes an irreversible phase transition from the pristine layered phase to a spinel-like phase and then to a completely inert rock salt phase in the deep delithiated state; this reconstruction process is accompanied by the rearrangement and release of lattice oxygen, as shown in Fig. 11g and h. Atomic-resolution electron energy loss spectroscopy (EELS) reveals the electronic structure of oxygen in different regions of the 4.8 V-charged sample. The pre-edge intensity of the O K-edge spectrum gradually diminishes from the material interior toward the surface, indicating an increase in oxygen vacancies and a decrease in oxygen atoms coordinated with TM ions.²²⁶ Meanwhile, as described in Section 2.2.2.2, cation mixing inherent to Ni-rich cathodes and exacerbated during charge/discharge processes can trigger electron rearrangement during deep delithiation. This process excites lattice oxygen activity, ultimately leading to lattice oxygen release, particularly in surface regions. This lattice oxygen release process, driven by the combined effects of weakened bond energy, lattice phase transitions, and cation mixing, ultimately manifests as a sustained degradation of both material structure and electrochemical performance.

It is noteworthy that the redox reaction of lattice oxygen requires sufficiently high charging voltages for activation. Typically, at low voltages, the capacity of Ni-rich cathodes is primarily contributed by the redox reaction of TM cations, with limited involvement of oxygen anion redox. At high voltages (4.5 V), both TM cation redox and oxygen anion redox contribute to capacity. However, this anion redox process is irreversible.^196,227 After lattice oxygen is released, it forms reactive oxygen species such as O₂²⁻ (peroxide), O₂⁻ (superoxide), O⁻, and molecular O₂. These highly reactive oxygen species readily react with trace H₂O and CO₂ present in the electrolyte, forming hydroxyl and carbonate species and releasing molecular O₂. Subsequently, these hydroxyl and carbonate species react with Li⁺ originating from the cathode, ultimately forming RLCs (mainly LiOH and Li₂CO₃). The reactions can be described as eqn (6)–(10).²²⁸ In addition, the Ni-rich cathode surface undergoes a transition from a layered structure to a rock salt phase after the release of lattice oxygen.^229,230 This cascade is regarded as a key trigger of thermal runaway.^{45,122,231,232} Thermally activated oxygen species engage in highly exothermic reactions with the organic electrolyte. The attendant heat evolution rapidly catalyzes electrolyte decomposition, undermines the cathode framework, and inexorably propels the cell toward catastrophic thermal runaway. Although the precise pathway of oxygen-evolution reactions in Ni-rich layered oxides remains elusive, the foregoing analysis indicates that oxygen release originates from lattice-oxygen oxidation. The reactive oxygen species produced in this process, together with continued lattice-oxygen depletion, drive a structural cascade from the layered phase to spinel and, ultimately, to a NiO-type rock salt framework.²³³ It is worth mentioning that oxygen release is inherently a bulk process, though its kinetics are partly limited by the long diffusion pathways of lattice oxygen. Thus, surface coatings, compositional-gradient designs, and Ni-concentration-engineered core–shell architectures are widely adopted to further curb oxygen evolution.

Ni³⁺ + O²⁻ (lattice) → Ni²⁺ + O⁻ (6)

O⁻ + O⁻ → O²⁻ (active) + O (7)

O²⁻ (active) + CO₂ → CO₃²⁻ (8)

O²⁻ (active) + H₂O → 2OH⁻ (9)

O + O → O₂ (10)

3.3. Electrolyte decomposition

Most of the gas generated is produced by the oxidative decomposition reaction of the electrolyte at high voltage which is divided into electrochemical oxidation and chemical oxidation. The electrolyte with ethyl carbonate (EC) as the solvent shows the following oxidative decomposition mechanism. Firstly, when the voltage is 3.8–4.0 V, the hydrogen atom on EC is attached to the oxygen atom of the layered oxide cathode.²³⁴ In the LiPF₆-based electrolyte, the active hydroxyl groups on the cathode surface could further react with the PF₆⁻ anion to release gaseous PF₅. PF₅ is extremely unstable and reacts with residual moisture in the electrolyte to form HF. Secondly, when the voltage is higher than 4.2 V (80% SOC), the chemical oxidation reaction mainly occurs, which depends on the release of reactive species. For example, in addition to the triplet O₂, the concentrated Ni⁴⁺ also induces the cathode to release singlet ¹O₂ from its layered structure. The singlet ¹O₂ tends to immediately nucleophilically attack EC to form a carbonyl group on EC, which is eventually decomposed into CO₂, CO and H₂O. CO is the concomitant product of CO₂, and the amount of CO is much lower than that of CO₂. The reaction is as shown in eqn (11):

2¹O₂ + EC → 2CO₂ + CO + 2H₂O (11)

Moreover, active proton-containing substances including acids and ethylene glycol, which originate from the hydrolysis of organic solvents and lithium salts, could react with LiPF₆, resulting in the formation of LiF, HF, HPO₂F₂ and POF₃ gases. Thirdly, under the overvoltage of >4.7 V, gas generation mainly comes from electrochemical oxidation. Electrochemical oxidation depends on the catalytic activity, specific surface area and voltage window of the cathode material. EC releases CO and CO₂ through a ring-opening reaction under the catalysis of the oxide cathode. In addition to the solvent, the carbon additives in the electrode could also be decomposed to release CO and CO₂. The whole reactions are represented in Fig. 11i.²³⁵

In addition, due to the presence of trace H₂O in the actual electrolyte, the H₂O in the electrolyte will decompose to produce H₂ when the decomposition potential of H₂O is reached (eqn (12)). According to the changes in ultrasonic transmission images during cycling of batteries assembled with electrolytes of different water contents (Fig. 11j), it can be seen that the higher the water content of the electrolyte, the earlier more gases appear inside the battery during cycling.²³⁶

2H₂O + 2e⁻ → H₂ + 2OH⁻ (12)

In fact, the water content in standard commercial electrolytes is very low. However, the amount of H₂ detected in batteries far exceeds the upper limit of the amount of H₂ that can be produced by the decomposition of trace amounts of water in the electrolyte. Metzger et al.²³⁷ found that a completely dry graphite/NCM full cell produces far more H₂ gas than that produced by the reduction of trace H₂O contaminants in the electrolyte and electrodes. They attributed this to the fact that solvents in the electrolyte oxidize to protonated solvents (denoted as R-H⁺) on the NCM surface, which subsequently diffuse to the anode surface to be reduced, especially at high temperatures and high voltages (Fig. 11k). The process is affected by the catalytic influence of TMs;²³⁸ therefore, Ni-rich cathodes tend to face more severe gas production phenomena due to the high Ni content and high catalytic activity of Ni.

4. Strategies to suppress the structural damage and gas emissions

To meet the rising demand for high-energy-density batteries, in-depth investigations have clarified the structural degradation and gas-evolution pathways of Ni-rich layered cathodes, thereby informing a suite of stabilization strategies. Common strategies include surface modification, element doping, electrolyte design and structural engineering.

4.1. Surface modification

Surface coatings act as a physical barrier that isolate the contact between the electrolyte and the Ni-rich cathode surface, thus suppressing parasitic interfacial reactions and micro-crack formation.^239–244 Appropriately engineered coating compositions could further stabilize surface lattice oxygen, inhibit gas evolution, and facilitate Li⁺ ion and electronic transport, thereby improving electrode kinetics. Some coatings directly convert RLCs into functional and well-ordered surface layers, markedly enhancing structural robustness. In addition, some surface coatings may form near-surface ionic doping synchronously during the coating process, further improving the structural stability of the material. In these coating strategies, coating materials are numerous, including metal oxide materials, lithium-ion conductors, and carbon materials.^38,245–249

Metal oxide coatings shield Ni-rich cathodes by physically separating them from the electrolyte and markedly mitigate particle fracture and cracking during repeated cycling. Currently, common metal oxide coating materials include Al₂O₃,²⁵⁰ WO₃,²⁵¹ ZrO₂^252,253 and others.^254,255 However, such surface coatings are generally electrochemically inert and have poor electrical conductivity. What's worse, overcoating usually results in a decrease in the rate performance and discharge capacity of the cathode material. As shown in Fig. 12a–c, Cao et al.²⁵⁶ employed a wet chemical method to coat NCM surfaces with varying amounts of Al₂O₃. Although this physical barrier enhanced cycling performance, the electrochemically inert Al₂O₃ reduced the discharge specific capacity of NCM. Moreover, due to low ionic conductivity, the coating exhibits poor enhancement effects on rate performance. When the coating content is excessive, it may even lead to a deterioration in rate performance. Therefore, when selecting a metal oxide for coating, the selection of dosage is very important. To achieve more precise coating coverage, Zhao et al.²⁵⁷ employed atomic layer deposition (ALD) to realize thinner and more uniform Al₂O₃ coatings. However, coating individual NCM secondary particles with inert metal oxides still inhibits ion and electron migration. Based on this, Wang et al.²⁵⁸ employed ALD to coat composite NCM811 cathodes with Al₂O₃ as shown in Fig. 12d (designated as AO-n, where n corresponds to the number of ALD cycles for thickness control of the Al₂O₃ coating layer). By comparing the electrochemical performance of NCM with varying coating thicknesses, they observed that due to the low lithium-ion conductivity of Al₂O₃, the specific capacity and coulombic efficiency of the AO-5, AO-10, and AO-50 samples were lower than those of the pristine NCM811 (Fig. 12e). However, the moderately coated AO-10 sample effectively suppressed side reactions between the cathode and the electrolyte while enhancing structural stability. This achieved optimal cycling stability alongside a significant improvement in rate performance (Fig. 12f). However, early studies on traditional inert metal oxide coatings often focused solely on the coating's ability to suppress interfacial side reactions and were oriented toward performance enhancement while neglecting the coating's impact on the crystal lattice and its inhibitory effect on oxygen release.

Fig. 12 (a) Al₂O₃ coating process on NCM cathodes. (b) Initial 0.1 C charge/discharge profiles and (c) rate capability of pristine and Al₂O₃ coated NCM samples.²⁵⁶ Copyright 2020, Elsevier. (d) Fabrication process and Li⁺ transport schematics comparing coating on powder vs. electrode using the ALD method. (e) Initial 0.1 C charge/discharge profiles and (f) rate performance of AO-n samples.²⁵⁸ Copyright 2022, Elsevier. (g) Schematic diagram of cation disorder of NCM and Ce_0.8Dy_0.2O_1.9-modified NCM (up) and schematic illustration of the oxygen vacancies in the Ce_0.8Dy_0.2O_1.9-modified layer enhancing the structural stability of Ni-rich materials (bottom). (h) Cycling performance at 1 C and (i) rate capability of pristine NCM and Ce_0.8Dy_0.2O_1.9-modified NCM. (j) Differential scanning calorimetry (DSC) traces showing the heat flow of NCM and NCM-CD3 charged to 4.3 V. (k) Ex situ X-ray diffraction (XRD) patterns of the pristine and NCM-CD3 electrodes after 100 cycles at 1 C.²⁶⁰ Copyright 2019, American Chemical Society. (l) Schematic diagram of effectively suppressing Li/Ni cation disorder via the nanoscale surface doping approach and reaction mechanisms of the oxygen ion conductor Ce_0.8Dy_0.2O_1.9 and lithium-ion conductor Li₈CeO₆ multifunctional surficial modification layer to enhance interfacial stability and safety performance of Ni-rich materials. (m) Contour plots of in situ XRD patterns of NCM (top) and NCM-LCD3 (bottom) for the (003) peak. (n) Voltage profiles of NCM and NCM-LCD3 at 2.8–4.7 V and the corresponding gas evolution during the first cycle with a current density of 200 mA g⁻¹ through in situ differential electrochemical mass spectroscopy measurements (DEMS). (o) Oxygen vacancy formation energy of NCM and oxygen ion conductor Ce_0.8Dy_0.2O_1.9 on the particle surfaces. (p) Comparison of long-term cycling performance of pristine NCM and NCM-LCD3.²⁶¹ Copyright 2023, John Wiley and Sons.

With a deepening understanding of the degradation mechanisms in Ni-rich cathodes and the demand to expand their practical applications, researchers have recognized that lattice oxygen release and the resulting chain reactions are the root causes limiting high-voltage applications of Ni-rich cathodes and leading to thermal runaway safety issues.²⁵⁹ Consequently, the goal of coating design has expanded beyond conventional performance metrics to focus on enhancing lattice oxygen stability while suppressing surface side reactions, thereby inhibiting oxygen release and structural degradation under high-voltage conditions. This transformation has given rise to a new generation of metal oxide functional coatings represented by oxygen ion conductors. Beyond providing conventional physical isolation, these materials actively guide and dissipate reactive oxygen species escaping from the lattice under high voltage through their internal oxygen vacancies and rapid oxygen ion migration pathways. This effectively suppresses the accumulation and release of oxygen, mitigating oxygen-induced structural degradation from the source while significantly enhancing material thermal stability under extreme conditions. Wang et al.²⁶⁰ coated NCM811 with oxygen ion conductor Ce_0.8Dy_0.2O_1.9 (NCM-CD) to achieve excellent electrochemical properties (Fig. 12g–i). The abundant oxygen vacancies in Ce_0.8Dy_0.2O_1.9 inhibit the release of oxygen, which improves the thermal stability and safety of Ni-rich cathode materials, as shown in Fig. 12j. Moreover, during the coating process, Ce⁴⁺ doping into the material lattice inhibits the disorder of Li⁺/Ni²⁺ cations and improves the structural stability of the material, enabling the cathode material to maintain an ordered structure after cycling (Fig. 12k). Benefiting from the suppressed interfacial side reactions, reduced irreversible surface structure degradation, and inhibited oxygen release enabled by this coating strategy, NCM-CD ultimately achieved satisfactory cycling and rate performance. Dai et al.²⁴⁰ proposed a kind of inert phase (La₂Mo₂O₉) with abundant oxygen vacancies, which is used as the surface coating of Ni-rich cathodes (L-NCM) and can act as both an oxygen anchor and an oxygen storage unit (Fig. 13a). With this modification scheme, the H2–H3 phase transition of the Ni-rich cathode is significantly suppressed (Fig. 13b), while the stability of lattice oxygen is enhanced due to the higher binding energies of interfacial La–O and Mo–O compared to Ni–O (Fig. 13c). In addition, the protective effect of the coating also leads to a weakened reaction of the electrolyte with oxygen. These favorable effects effectively inhibit the structural deterioration and improve the electrochemical performance of the Ni-rich cathode, resulting in a capacity retention rate of 94% for the full cell after 400 cycles (Fig. 13d).

Fig. 13 (a) Schematic illustrations of the operando oxygen anchoring and reserving strategy. (b) Selected (003) counter plot of in situ XRD patterns for pristine NCM (P-NCM) and L-NCM upon initial cycle. (c) Projected density of states (DOS) for surface oxygen coordinated by one La, two Li and two Ni (left), and surface oxygen coordinated by one Mo, two Li and two Ni (right). (d) Cycling performance of the L-NCM/Gr pouch-type full cell at 0.5 C under 2.7–4.2 V. Inset: Photo of the assembled pouch cell (size: 5 cm long and 8 cm wide).²⁴⁰ Copyright 2023, Springer Nature.

To simultaneously stabilize the cathode lattice, accelerate Li⁺ transport and suppress R_ct, lithium-ion-conductive surface coatings were used to modify the Ni-rich cathodes.^262,263 Common lithium-ion conductor coatings include Li₄Ti₅O₁₂,²⁶⁴ Li₂WO₄,²⁶⁵ and so on. Tang et al.²⁶⁶ used NASICON-type Li_1.3Y_0.3Zr_1.7(PO₄)₃ (LYZP) to coat single-crystal LiNi_0.83Co_0.11Mn_0.06O₂ (SC-NCM83) to form SC-NCM83@Zr/LYZP, and introduced trace amounts of Zr doping in the formation of a coating layer (Fig. 14a). In situ XRD tests show that SC-NCM83@Zr/LYZP displays a reversible and weaker H2–H3 phase transition with a maximum c-axis shrinkage of 2.60%, an a-axis change of 1.88%, and a volume change of 4.89%. The lattice expansion/contraction during the first charge–discharge cycle is significantly lower than that of the pristine SC-NCM811 (Fig. 14b, c and g), enhancing the integrity of the layered structure. In addition, SC-NCM83@Zr/LYZP achieves more stable lattice oxygen and suppressed surface side reactions with significantly lower O₂ and CO₂ release, as shown in Fig. 14d and e. Compared with SC-NCM83, the exothermic peak of SC-NCM83@Zr/LYZP shifts backward and the exothermic amount is significantly reduced (Fig. 14h), demonstrating excellent thermal stability. Therefore, the capacity retention of SC-NCM83@Zr/LYZP in the extended voltage range of 2.75–4.5 V is 81.2%, and it also exhibits higher stability at a high current density of 1 C, retaining 77.3% capacity compared to 65.6% for the unmodified SC-NCM83 (Fig. 14f). Wang et al.²⁶⁷ coated LiNi_0.9Co_0.05Mn_0.05O₂ (NCM9) with the perovskite lithium/oxygen dual-ion conductor La₄NiLiO₈ (LNLO) (Fig. 14i). The stable oxygen vacancies/interstitials in LNLO can effectively stabilize the lattice oxygen on the surface of the Ni-rich cathode. The gas production problem of the Ni-rich cathode is fundamentally reduced (Fig. 14j), and the change in the lattice constant during cycling is increased as well, which improves phase transition reversibility and inhibits particle breakup after long cycling (Fig. 14k). As a result, the material shows excellent cycling stability, and the capacity retention rate of 66.33% is still achieved at 400 cycles at a high voltage of 4.5 V (Fig. 14l). At the same time, the good lithium-ion conductivity also allows a better rate performance. Moreover, on the basis of Ce_0.8Dy_0.2O_1.9-coated NCM811, Wang et al.²⁶⁸ further designed a modification strategy based on combining the metal oxide Ce_0.8Dy_0.2O_1.9 with the lithium-ion conductor Li₈CeO₆ as the surface coating (NCM-LCD3) as shown in Fig. 12l, which suppressed the H2–H3 phase transition (Fig. 12m) and reduced the deleterious gas production phenomenon (Fig. 12n). The inhibition of oxygen release is attributed to the high oxygen vacancy formation energy of Ce_0.8Dy_0.2O_1.9 as shown in Fig. 12o. As a result, NCM-LCD3 demonstrated excellent high voltage cycling stability with a cycle life of over 400 cycles (Fig. 12p).

Fig. 14 (a) Schematic of the SC-NCM83@Zr/LYZP synthesis process. In situ XRD tests of (b) SC-NCM83@Zr/LYZP2 and (c) pristine SC-NCM83 during initial cycling at 0.1 C between 2.75 and 4.5 V. The corresponding variations in the (g) c-axis parameter, a-axis parameter, and cell volume during the charging process. In situ DEMS data of (d) SC-NCM83@Zr/LYZP2 and (e) SC-NCM83 measured under the same electrochemical conditions as in situ XRD. (h) DSC curves of SC-NCM83@Zr/LYZP2 and SC-NCM83 electrodes after charging to 4.5 V. (f) Cycling performance of SC-NCM83@Zr/LYZP2 and SC-NCM83 at different current density. (h) DSC curves of SC-NCM83@Zr/LYZP2 and SC-NCM83 electrodes after charging to 4.5 V. ²⁶⁶ Copyright 2025, Elsevier. (i) Schematic diagram of effectively inhibiting the release of highly active oxygenates in the surficial lattice by oxygen vacancies and oxygen interstitials in the layered perovskite LNLO to improve battery safety performance and alleviate stability issues of ultrahigh-Ni layered oxide cathode material NCM9 (NCM9-LNLO). (j) Voltage profiles of as-prepared electrodes at 2.8–4.7 V and the corresponding gas evolution during the first cycle of NCM9 and NCM9-LNLO3 with a current density of 200 mA g⁻¹ through in situ DEMS measurements. (k) Comparison of the changes of lattice parameter c obtained by fitting XRD patterns during the first two charge/discharge processes of NCM9 and NCM9-LNLO3. (l) Comparison of the long-term cycling performance of pristine NCM9 and NCM9-LNLO3 at 2.8–4.5 V.²⁶⁷ Copyright 2023, John Wiley and Sons.

The carbon material could improve the electron and ion conductivity, prevent the structural collapse caused by the dissolution of TM ions, and effectively suppress the capacity decay. Guo et al.²⁶⁹ prepared carbon-coated LiNi_0.92Co_0.06Al_0.02O₂ nanosheets by sol–gel assisted high-temperature annealing. The carbon coating can protect the electrode from the influence of H₂O and CO₂, which enhances the electrode's corrosion resistance and makes it easier to store. The modified cathode exhibits excellent structural integrity, and the battery capacity is also improved. Liu et al.²⁷⁰ used sucrose and glucose as carbon materials to construct a nano-carbon coating on the surface. It is found that carbon coating, as a physical barrier between the cathode material and the electrolyte, can effectively prevent the formation of microcracks and inhibit the side reaction at the interface, and improve the rate and cycling performance of the cathode. In addition, the coating effect of sucrose is better than that of glucose. The reason is that the glucose coating is dense, while the sucrose coating is relatively loose and has a larger specific surface area. Therefore, the sucrose coating can inhibit the erosion of the electrolyte more effectively.

In addition to constructing coating layers on the surface, researchers have devised a variety of novel surface treatments in recent years. Sun et al.⁵⁶ proposed a washing process using an aqueous solution of Co that simultaneously removes RLCs and forms a dual protective coating of an outer F-rich layer and an inner Co-rich layer on the Ni-rich layered cathode (Fig. 15a and b). The Co-rich layer formed during the washing process depletes the surface RLCs and prevents direct contact of the electrolyte with the surface of the Ni-rich cathode, improving the lattice oxygen stability. The additional fluorine coating further reacts with the surface RLCs to generate a prefabricated LiF coating (Fig. 15c), which effectively mitigates the decomposition of electrolytes and impurities, improves the interfacial stability, and suppresses gas generation (Fig. 15d). The synergistic effect of the two modifications together improves the long-term cycling stability of the Ni-rich cathode (Fig. 15e). Cai et al.²⁷¹ proposed a surface lanthanide strategy for stabilizing the high-voltage cathode based on the general principles of high-voltage cathode near-surface design using LCO as a substrate, which modulates the near-surface structure of the material (La-LCO). An implanted La/Ca gradient is formed by the exchange of La³⁺/Ca²⁺ with Li⁺ in aqueous solution. After the annealing treatment, a perovskite protective layer La_1−wCa_wCoO_3−δ was formed on the surface of LCO, while the near surface showed a gradient doping of La³⁺/Ca²⁺ (Fig. 15f). The nanoscale La_1−wCa_wCoO_3−δ surface perovskite phase functions as an oxygen buffer that reversibly stores oxygen while elevating the lattice-oxygen release potential, as proved by DEMS tests, thereby effectively suppressing oxygen and other derivative gas release. After 100 cycles at 1 C and 3.0–4.5 V, the pouch cell of La-LCO exhibits nearly imperceptible swelling, in stark contrast to the pronounced bulging observed in the pristine LCO (P-LCO) pouch cell (Fig. 15g and h). Concurrently, it mitigates electrochemical degradation caused by interfacial side reactions and structural decay. And it was generalized to Ni-rich layered cathodes (La-NCM) with high energy density (Fig. 15i). In order to solve the limitation that the current coating means are only limited to the secondary particle surface coating but cannot inhibit the side reactions at the root of the microscopic solid–liquid interface between primary particles, Yoon et al.²⁴² chose to select a cobalt boride (Co_xB) for the simultaneous coating of both the primary particle interface and the secondary particle surface of NCM811 (Fig. 15j). There is a strong reactivity between Co_xB and the surface oxygen of NCM811 and a passivation film can be formed without consuming too much lattice oxygen. Eventually, primary particle cracking and surface transition were suppressed (Fig. 15k), and gas escape was reduced (Fig. 15l), thus effectively improving rate performance and high temperature cycling stability (Fig. 15m and n). Sun et al.²⁷² applied thermal quenching (140 °C min⁻¹) to high-Ni LiNi_xMn_yCo_zO₂ (x ≥ 0.8) immediately after calcination. The non-equilibrium cooling expels Li⁺ toward the surface, creating a nano-crystalline Li₂CO₃ overlayer that passivates against electrolyte oxidation and stabilizes lattice oxygen. Quenched NMC-Q exhibits 195 mAh g⁻¹ at 100 mA g⁻¹ and retains 90% capacity after 300 cycles at 200 mA g⁻¹. Wang et al.²⁷³ performed rapid deionized water treatment on NCM811 followed by subsequent heat treatment. The treatment leaches a few surface Li⁺ layers and triggers in-plane Ni–O rearrangement, yielding a 2–4 nm rocksalt NiO-like phase that is coherent with the underlying layered lattice.This surface reconstruction phase passivates the material surface and forms stable self-anchored Ni–O clusters, thereby suppressing electrochemical cycling degradation in Ni-rich materials. The modified NCM maintains 86% retention after 300 cycles at 1 C without external dopants or coatings. Lu et al.²⁷⁴ developed a surface solid reaction strategy to stabilize LiNi_0.83Co_0.07Mn_0.1O₂ cathodes by constructing a conformal La(OH)₃ nanocoating via an anhydrous wet-chemical method. Using ethanol and hexamethylenetediamine, this decomposition-induced precipitation technique achieved nanoscale-precision coating without damaging the moisture-sensitive structure. Following high-temperature sintering, this coating undergoes a controlled solid reaction with the NCM surface. This process introduces moderate Li/Ni disorder into the cathode surface layer, which synergistically interacts with the generated LiLaO₂ to enhance the cathode's cycling performance. After 400 cycles at 1 C, capacity retention jumps from 35.5% to 80.8%, while 5 C discharge remains at 144 mAh g⁻¹. Notably, the anhydrous medium avoids proton or carbonate contamination that typically plagues aqueous coating processes.

Fig. 15 (a) and (b) A TEM image and the corresponding energy dispersive spectrometer (EDS) overlay map showing Co and F distribution (a) and EDS line scan results of the surface of the F-Co-washed cathode (b) along a green arrow in panel a. The vertical dashed lines in panel b indicate the thickness of each coating layer. (c) Comparison of F 1s, O 1s and Li 1s X-ray photoelectron spectroscopy (XPS) spectra of the unwashed, F-coated, deionized water-washed (DI-washed), Co-washed and F-Co-washed cathodes. The vertical dashed lines in panel c indicate the binding energy of XPS peaks for each cathode, and the dashed curves indicate the fitted profiles by peaks from the residual lithium compounds and intercalated lithium on the surface. (d) Gas evolution from pouches featuring a 4.3 V charged cathode and the electrolyte during storage at 60 °C. In panel (d) black filled-square data represent the volume changes of the pouch containing the electrolyte and the uncharged unwashed cathode. (e) Long-term cycling performance using pouch-type full cells.⁵⁶ Copyright 2023, Springer Nature. (f) Schematic of the treatment process, the resulting surface architecture with a La/Ca gradient and perovskite layer, and the corresponding chemical potential (µ) profile of La³⁺/Ca²⁺. (g) Optical photos of P-LCO and La-LCO pouch cells after 100 cycles at 1 C (3.0–4.5 V), demonstrating the enhanced physical integrity. (h) In situ DEMS profiles during the first charge to 4.7 V, confirming the suppression of oxygen-related gas release in the La-LCO half-cell. (i) Cycling performance of P-NCM and La-NCM in coin-type half-cells at 1 C at 2.8–4.4 V vs. Li⁺/Li.²⁷¹ Copyright 2023, Springer Nature. (j) Schematic coating-plus-infusion microstructure in which Co_xB uniformly coats the surface of NCM secondary particles and infuses into the grain boundaries between the NCM primary particles. (k) TEM image and EDS mapping (inset) of Co_xB-NCM after 200 cycles at a 7 C discharge rate at 45 °C and HR-TEM of a secondary particle surface in cycled Co_xB-NCM. (l) In situ DEMS data of pristine NCM (top) and Co_xB-NCM (bottom) during the first charge at 0.2 C in the voltage range of 3.0–4.4 V versus Li/Li⁺ at 25 °C. (m) Rate tests and 1 C cycling of Co_xB-NCM and pristine NCM in the range of 3.0–4.4 V versus Li/Li⁺ at 25 °C. The shading shows the standard deviations calculated from five cells. (n) 7 C discharge cycling tests in the range of 3.0–4.4 V versus Li/Li⁺ at 45 °C, with 6 intermittent cycles with 0.2 C charge/discharge conducted after every 100 cycles.²⁴² Copyright 2021, Springer Nature.

4.2. Element doping

Element doping is also an effective strategy to improve the electrochemical performance of Ni-rich layered oxide cathodes.^275–277 This simple modification improves the stability of the lattice structure at the atomic level, increases the bonding strength between the metal ions and the oxygen atoms, which reduces the degree of cation mixing and irreversible phase transitions, and prevents microcracking. In addition, proper doping can promote the migration of Li⁺ and refine the microstructure of the material.^278–281 So far, it has been reported that strategies such as single cation doping, single anion doping and mixed ion doping can be used to improve the comprehensive properties of Ni-rich cathode materials.

Cation doping can effectively improve the cycling performance and internal structure stability of Ni-rich cathodes.^224,282 The different doping sites depend on the radius of the doped cation. Some small radius cations, such as K⁺, Na⁺, and Mg²⁺, are more inclined to occupy the position of lithium ions, inhibit cation mixing and play a columnar effect. The columnar effect can significantly inhibit the formation and expansion of microcracks, so as to obtain better cyclic performance.²⁸³ However, this type of doping may also lead to capacity degradation. Using a simple mixed molten salt sintering process, Jian et al.²⁸⁴ synthesized a Na-doped single crystal LiNi_0.82Co_0.125Mn_0.055O₂ (NCM-Na) Ni-rich cathode. In this case, Na⁺ uniformly occupies the Li sites, and Na⁺ can selectively stabilize the surrounding Li⁺ by regulating the orientation of the Jahn-Teller effect of Ni³⁺. These stabilized Li ions provide preferential channels that enable high-speed movement of Li ions. Benefiting from this, the Li inhomogeneity and phase separation during charging and discharging of NCM-Na are greatly reduced, thus reducing the risk of structural defects and mechanical degradation and improving cycling stability and rate performance. Qiu et al.²⁸⁵ pioneered a novel doping process and proposed the concept of a Mg²⁺ releasing membrane for the first time. That is, the doping of Mg²⁺ is realized during the cycling process. A self-supporting MgV₂O₄ membrane was synthesized by incorporating mobile Mg²⁺ into the MgV₂O₄ structure, which releases Mg²⁺ during electrochemical processes and acts as a “pillar” within the lattice to provide structural reinforcement to the cathode surface.

Cations with large ionic radius, such as Ti⁴⁺, Ta⁵⁺, Mo⁶⁺, etc., often occupy the position of the TM layer and enhance the strength of the metal–O bond in the TM layer, and thus improve the lattice oxygen stability of the Ni-rich cathode.⁸¹ Park et al.²⁸⁶ systematically investigated the effect of doping of various elements, including Co, Al, Ti, Ni, Ta, W, Mo, etc., on the LiNi_0.9Mn_0.1O₂ (NM90) cathode. It was found that doping of high valence elements could not only inhibit the H2–H3 phase transition more effectively but also refine the primary grain size (Fig. 16a). As a result, better cycling performance is presented (Fig. 16b and c). After matching the fluorine-containing electrolyte with Mo-doped NM90, the capacity retention rate of the full cell reached 85.6% after 1000 cycles at 4.3 V. Further, Ryu et al.²⁸⁷ studied in detail the effect of Mo doping amount and synthesis temperature on a Ni/Co/Mn/Al quaternary Ni-rich cathode (NCMA94) containing 94% Ni. The results show that the electrochemical properties of the Mo doped Ni-rich layered cathode are largely determined by the primary particle size and morphology of the cathode material and the crystallinity of the material. The higher the crystallinity and the finer the primary particles, the better the capacity and cycling performance of the Ni-rich cathode. In addition, the primary particles become finer with the increase of Mo content (Fig. 16d), while the strength of the material tends to stabilize at 1% Mo content (Fig. 16e). The electrochemical performance of the Ni-rich cathode is improved after Mo element doping (Fig. 16f). Yang et al.²²⁴ constructed LiNi_0.94Co_0.05Te_0.01O₂ (NC95T) using Te⁶⁺ doping, and the introduction of Te⁶⁺ led to the formation of a highly thermodynamically stable Ni–Te ordered structure within the TM layer and improved particle morphology. This micromorphology and ordered atomic arrangement resulted in a reduction of lattice strain during cycling for NC95T (Fig. 16g and h), and significantly improved the lattice oxygen stability. Further calculations show that for the NC95T cathode, especially in the fully delithiated state, a localized O 2p state near the oxygen Fermi energy level appears in the ordered Ni–Te structure, which makes the electrons in the lattice oxygen more stable (Fig. 16i). The Bader charge of the NC95T cathode has a discrete nature, which suggests the passivation of the lattice oxygen (Fig. 16j). In addition, the introduction of the ordered Ni–Te structure results in a high oxygen vacancy formation energy for NC95T (Fig. 16k), which also confirms the excellent stability of lattice oxygen.

Fig. 16 (a) Comparison of primary particle morphologies of undoped and doped NM90 cathodes. (b) and (c) Comparison of NM90 cycling performance after doping with (b) low valence cations and (c) high valence cations.²⁸⁶ Copyright 2022, Springer Nature. (d) Distributions of the primary particle size. (e) Secondary particle strength of the optimized NCMA94 and Mo-doped NCMA94 cathodes. (f) Long-term cycling performance of optimized NCMA94, Mo1-NCMA94, and Mo1.5-NCMA cathodes in full cells.²⁸⁷ Copyright 2023, Elsevier. (g) In situ XRD of NC95T of the first two cycles and (h) corresponding calculated lattice strain. (i) The density of states of oxygen. (j) Bader charge on oxygen and average Bader charge difference between the pristine and Ni–Te ordered sample. (k) Oxygen vacancy formation energy (pristine means NC95 and Ni–Te ordered represents NC95T).²²⁴ Copyright 2024, Springer Nature.

Anion doping, when compared with cation doping, is suggested to provide stronger solubility resistance in acid for Ni-rich materials.²⁸⁸ It can also prevent Li or TM ions from occupying incorrect sites and improve the energy density of the electrode.²⁸⁹ When the oxygen atom in the layered cathode part is replaced by a halogen atom (X), the formed X-TM (Ni, Co, Mn) has a higher bonding strength than the original O–TM (Ni, Co, Mn), thus enhancing the structural strength of the Ni-rich cathode. Zhang et al.²⁹⁰ found that the F⁻ doping strategy can reduce the Li/Ni mixing, and accelerating rate calorimetry studies suggest that samples with higher F content in the cathode will have safety advantages. Similarly, F-doped NCM811 was prepared by a solid-state reaction method. The modified samples show better cycling performance due to the strong bonding of TM and F as well as improved Li⁺ transport behavior.²⁹¹ In addition, Park et al.²⁹² found that the capacity retention of sulfur-doped LiNiO₂ cathode materials was significantly improved. Since the electronegativity of sulfur is lower than that of oxygen, sulfur doping can prevent interlayer elongation and structural decomposition caused by lithiation/delithiation during the cycle by increasing the transport of Li⁺. Cl⁻ doping has a good inhibitory effect on Li/Ni disorder and could increase the diffusion rate of Li⁺ as well.²⁸⁹

To combine the advantages of different ions, some researchers have adopted the strategy of mixed ion doping. The synergistic effect between dopants is utilized to maximize the performance of Ni-rich cathodes. Lee et al.⁸⁰ improved the stability of the LiNi_0.92Co_0.04Mn_0.04O₂ (NCM92) cathode based on a double doping strategy by employing Al³⁺ and Nb⁵⁺ (AlNb-NCM92). Al³⁺ doping enhances the stability of the crystal structure, while Nb⁵⁺ doping optimizes the morphology of primary particles. Under the synergistic action of the two elements, the lattice ordering of the Ni-rich cathode increases and the lattice strain decreases (Fig. 17a–c). Microcracks are almost absent in AlNb-NCM92 after 1000 cycles, which is attributed to the fine and radially aligned primary particles effectively dispersing the strains imposed on the particles during the charge/discharge process, thus suppressing the generation of microcracks (Fig. 17e–g). The cathode can still maintain 88.3% of the initial capacity after 1000 cycles. Zhang et al.²⁹³ found through experiments that Cr–Mg co-doping can lead to a synergistic reaction to form a complementary structure. Cr occupies the Ni layer and Mg occupies the Li layer. The synergistic reaction can reduce the cation mixing degree and improve the stability of the structure, thus enhancing the cycle performance of the cathode and showing a higher discharge capacity. Zhou et al.²⁹⁴ mitigated the structural degradation and oxygen release problems of Ni-rich cathode after prolonged cycling by employing Mg²⁺/F⁻ co-doping (Mg1+F2) (Fig. 17d). The co-doping strategy results in a high Ni²⁺ migration barrier and high oxygen vacancy formation energy (Fig. 17h–j). In situ XRD tests reveal a more stable and reversible structure of Mg1+F2, which exhibits lower Li/Ni mixing and suppressed H2–H3 phase transition (Fig. 17k–m and p). Electron paramagnetic resonance (EPR) results indicate (Fig. 17n and o) that the oxygen vacancy concentration of Mg1+F2 in the charged state is almost unchanged compared with that in the pristine state, whereas the oxygen vacancy concentration of the unmodified Ni-rich cathode increased significantly. The results of in situ DEMS show that the CO₂ release from Mg1+F2 is reduced, and the onset voltage of the release is delayed, indicating that the cathode/electrolyte interfacial side reactions are suppressed. Notably, there is no O₂ release observed in the test, indicating that the co-doping of Mg²⁺ and F⁻ can effectively stabilize the lattice oxygen and thus inhibit the O₂ release (Fig. 17q and r). For LiNi_0.96Co_0.03Mn_0.01O₂, Xu et al.²⁹⁵ combined trace W-doping (microstructural refinement) with a conformal LiF layer (grain-boundary sealing). The W-doped matrix forms slender primary particles that relieve inter-granular stress; the LiF layer blocks HF attack and electrolyte penetration.

Fig. 17 Magnified XRD patterns highlighting (003) and (104) peaks (a) and (b) cation mixing of Ni²⁺ in the Li layer of the as-prepared NCM92, Al-NCM92, Nb-NCM92 and AlNb-NCM92. (c) Corresponding normalized lattice volume variations of the cathodes measured by in situ XRD.⁸⁰ Copyright 2024, American Chemical Society. (d) Schematic diagram of the structural evolution of NCA and Mg1+F2.²⁹⁴ Copyright 2023, American Chemical Society. Cross-sectional scanning electron microscope (SEM) images of (e) discharged NCM92, Al-NCM92, Nb-NCM92 and AlNb-NCM92 cathode particles and charged (f) Nb-NCM92 and (g) AlNb-NCM92 cathode particles after long-term cycles.⁸⁰ Copyright 2024, American Chemical Society. (h) Simplified model of Ni migration to the Li site and simplified model of lattice oxygen escape. (i) Migration barrier of Ni²⁺ in NCA and Mg1+F2 to the Li layer. (j) Escape barrier of lattice oxygen in NCA and Mg1+F2. In situ XRD patterns of the (003) peak in the initial cycle for (k) NCA and (l) Mg1+F2. (m) Evolution of the (003)/(104) ratio for NCA and Mg1+F2. EPR of NCA and Mg1+F2 before (n) and after (o) cycling. (p) Evolution of lattice parameter c for NCA and Mg1 F2. In situ DEMS measurements for (q) NCA and (r) Mg1+F2.²⁹⁴ Copyright 2023, American Chemical Society.

4.3. Electrolyte design

Electrolyte engineering is one of the strategies that can significantly and simultaneously improve the cycle life and safety of Ni-rich cathodes. In high voltage cycles, the electrolyte would oxidize and decompose, corrode the surface of the cathode and form a CEI layer on the surface of the material, resulting in increased interface impedance and thus affecting the cycle performance of the battery.^{245,248,296–299} In addition to surface engineering, appropriate electrolyte additives could also be selected to improve the high voltage and high temperature stability of the electrolyte, thereby reducing the damage from the interface side reaction to the cathode material and forming a stable CEI on its surface.^{88,243,245,300–304} Common additives can be divided into solvents and lithium salts,^{296,305–318} such as fluoroethylene carbonate (FEC), vinyl carbonate (VC), LiNO₃, LiBF₄ and so on. In addition to a single additive, the simultaneous use of multiple additives can also produce synergistic effects. Zhang et al.³¹⁹ proposed a trimethyl borate (TMB) electrolyte additive with low oxidation potential that forms a stable and dense CEI, suppresses the dissolution of TMs, and eventually enables stable cycling of NCM90/Li batteries at high voltages of 4.7 V and 4.5 V. The NCM90/Li battery with 1 wt% TMB addition has a capacity retention rate of 77.8% after 200 cycles at 4.5 V and can still reach 70.7% after 100 cycles at an ultra-high voltage of 4.7 V (Fig. 18a–c). Zou et al.³²⁰ inhibited EC dehydrogenation on the cathode surface by introducing the 4-fluoro-N,N-dimethylbenzenesulfonamide electrolyte additive (FBSN) and formed a dense, rigid substance with high ionic conductivity and a self-healing CEI on the Ni-rich cathode surface. Therefore, the cycle stability of the Ni-rich cathode is obviously improved. Cheng et al. proposed to promote preferential deposition of anions by constructing dissolved sheaths enriched with multiple anions and to form solid electrolyte interfaces enriched with inorganics. Therefore, the performance of the Ni-rich cathode under high voltage and high temperature is effectively improved. At the same time, the existence of the weak solvation structure greatly reduces the R_ct at low temperature and inhibits the growth of lithium dendrites.^{297,321–323} This high-chaos electrolyte (HC-E) design strategy enables NCM811/Li to exhibit excellent electrochemical performance in the temperature range of −30 to 45 °C at 4.7 V(Fig. 18d–f).³²⁴ Chen et al.³²⁵ proposed a strategy involving the engineering of an insoluble CEI to address interface-induced degradation in Ni-rich cathodes. A series of siloxanes containing unsaturated units (propargyloxytrimethylsilane (PMSL), allyloxytrimethylsilane (AMSL), and trimethylmethoxysilane (TMSL)) were used as electrolyte additives (Fig. 18g). These siloxane compounds, featuring high unsaturation with numerous reactive sites, undergo increased polymerization reactions via 3D topological pathways, forming insoluble CEI materials. Thus, the side reactions between the cathode and the electrolyte, transition metal dissolution/crosstalk effects, as well as stress corrosion cracking and impedance increase in the active material are suppressed, as shown in Fig. 18h. Owing to the stable CEI, the NCM811/graphite pouch cell maintains a capacity retention rate exceeding 85% after approximately 300 cycles at 60 °C. The higher the unsaturation level of unsaturated units in the electrolyte additives, the lower the gas evolution in the pouch cell after prolonged cycling as shown in Fig. 18i.

Fig. 18 (a) Schematic diagram of the action mechanism of TMB on the NCM90 cathode surface. (b) and (c) Electrochemical performance of NCM90/Li cells with/without TMB.³¹⁹ Copyright 2024, Elsevier. Cycling performance of the NCM811/Li cells using HC-E and baseline electrolytes at (d) 45 °C and (e) −30 °C, and (f) first discharge specific capacity at different temperatures.³²⁴ Copyright 2023, American Chemical Society. (g) Molecular structures of TMSL, AMSL, and PMSL (with increasing unsaturation) and their proposed function in forming a passivation layer. (h) Origin of multifunctionality illustrating the stabilization of both electrode interphase and bulk electrolyte via possible electrochemical polymerization pathways. (i) Enhanced cycling stability of commercial NCM811/graphite pouch cells in base and additive-containing electrolytes at elevated temperatures.³²⁵ Copyright 2022, John Wiley and Sons.

With the shift in electrolyte design from relying on single functional components to developing synergistic electrolyte systems, locally high-concentration electrolytes (LHCEs) have emerged. LHCEs successfully addresses the bottlenecks of traditional high-concentration electrolytes (high viscosity, poor wetting, and high cost) by introducing inert diluents. This approach preserves the inherent advantages of high-concentration lithium salts, such as a rich inorganic interface and suppressed solvent side reactions, while effectively enhancing the interface stability of Ni-rich cathodes. For example, Jia et al.³²⁶ designed a non-flammable LHCE based on a mixed solvent of tetramethylene sulfone and trimethyl phosphate, combined with the 1H, 1H, 5H-octafluoropentyl 1,1,2,2-tetrafluoroethyl ether high flash point diluent (SEO-M1). Mechanistic studies indicate that compared to the baseline electrolyte (E-Baseline), the LHCE effectively mitigates irreversible phase transformations on the surface of Ni-rich cathodes. As shown in Fig. 19a and b, after cycling in the baseline electrolyte, NMC811 exhibits two thicker degradation phases: the rock salt phase and the cation mixed phase. In contrast, NMC811 cycled in SEO-M1 formed only a thin cation-mixed layer. Furthermore, NMC811 cycled in SEO-M1 exhibits a more inorganic-rich CEI with less organic carbonate and more LiF, confirming that this LHCE suppresses surface side reactions on NCM811 (Fig. 19c). Consequently, the optimal LHCE significantly enhances the cycling stability of graphite/NCM811 batteries after 500 cycles compared to the baseline electrolyte (Fig. 19d). Li et al.³²⁷ employed multifunctional flame-retardant diluent ethoxy (pentafluoro) cyclopropanecarbonitrile (PFPN) to regulate a locally high-concentration electrolyte (LHCE-PFPN) based on triethyl phosphate (TEP). This approach achieved outstanding flame retardancy while maintaining excellent compatibility with high-voltage cathodes and lithium metal anodes. PFPN preferentially reduces on the lithium metal surface to form LiF and PFPN derivatives. Subsequently, PFPN derivatives migrate to the NCM811 surface, undergoing ring-opening polymerization or oxidative decomposition to generate inorganic compounds, ultimately concurrently constructing an inorganic-rich SEI and CEI as shown in Fig. 19e. This stable interfacial structure maintains lithium metal anode stability while suppressing electrolyte decomposition on the NCM811 surface and reducing TM dissolution (Fig. 19f). Consequently, it significantly enhances the cycling performance of Li/NCM811 batteries under harsh high-temperature and high-voltage conditions, as demonstrated in Fig. 19g and h.

Fig. 19 (a) HAADF and (b) bright field STEM images of pristine NCM811, and after cycling with E-baseline and SEO-M1 at 45 °C, showing suppressed cation mixing and microcracking. (c) XPS spectra of C 1s, O 1s, and F 1s for graphite anodes cycled with E-Baseline and SEO-M1, indicating a superior CEI composition. (d) Average specific discharge capacity and coulombic efficiency of Gr/NMC811 full cells with different electrolytes at 25 °C and C/3 rate.³²⁶ Copyright 2025, American Chemical Society. (e) Reaction mechanism of PFPN in forming a robust SEI/CEI. (f) Schematic diagram comparing the solvation structures, interphase formation, and Li deposition behaviors in LHCE-PFPN and LHCE-1,1,2,2-tetrafluoroethyl-2,2,3,3-tetrafluoropropylether (TTE). (g) and (h) Cycle performance of Li/NCM811 cells demonstrating the superiority of the LHCE-PFPN electrolyte at 60 °C (g) and under a high cut-off voltage of 4.5 V (h).³²⁷ Copyright 2024, Elsevier.

4.4. Structural engineering

Structural engineering plays a key role in improving the structural stability of Ni-rich layered cathode materials.^{56,92,328–330} By precisely controlling the microstructure, structural damage and gas emission during charging and discharging can be effectively controlled.³³¹ Common structural engineering approaches include single-crystallization and concentration gradient design.^{275,332–334}

The secondary particles of polycrystalline Ni-rich cathode materials are aggregated from primary particles at the nanometer scale, with grain boundaries between the primary particles. Due to lithiation/delithiation during the charge/discharge process, the crystals continuously undergo anisotropic contraction/expansion, resulting in a high volume change rate and stress accumulation, which ultimately leads to contact failure between the primary particles and consequently rupture of the secondary particles. The electrolyte penetrates the material along the cracks eventually.^70,335 In contrast, single-crystal Ni-rich cathode particles do not have grain boundaries and thus can inhibit the penetration of the electrolyte. Moreover, single-crystal particles can eliminate primary intergranular stresses and effectively reduce the drastic changes in lattice parameters during the cycling process, thus reducing intergranular cracks and particle cracking and comminution caused by irreversible phase transitions (Fig. 20a).³³⁶ The single-crystal structure also has a high bulk density and is non-porous to withstand greater pressure.³³⁷ Nam et al.³³⁸ synthesized single-crystal Ni-rich (DSNCM) cathodes using a pellet-assisted mechanical densification process. As shown in Fig. 20b and c, after 150 cycles at high voltage, the single-crystal material can still maintain 85% capacity retention which is higher than that of the poly-crystal sample (PNCM). Moreover, the formation of intergranular cracks is inhibited, which is helpful to improve the mechanical stability of single-crystal materials. Liu et al.³³⁹ synthesized Ni-rich single-crystal materials using a binary molten salt method. Compared to the poly-crystal sample, the single-crystal sample shows a lower degree of cation mixing and larger lithium layer spacing, and the RLC is effectively removed from the surface of single-crystal particles.

Fig. 20 (a) Schematic diagram of structural changes during the cycle of single-crystal and poly-crystal Ni-rich cathodes.³³⁶ Copyright 2020, Elsevier. (b) Comparison of electrochemical properties of DSNCN and PNCM. (c) SEM of DSNCM and PNCM after cycling.³³⁸ Copyright 2023, Elsevier. (d) Rate capabilities and (e) chemical diffusion coefficients of lithium ions in P-NCM90 and three S-NCM cathodes. (f) TEM image of the S-NCM90 cathode particle charged to 4.3 V at 0.5 C and part of electron diffraction patterns from the regions marked (i) and (ii). The yellow arrows are parallel to the lithium-ion diffusion path in the particle. The dot-line diagram displays the c-axis lattice parameters at locations along the dashed yellow line. Overlays and deconvolution of the (003) reflection peaks recorded by in situ XRD in the voltage range of 4.15–4.5 V: (g) and (k) P-NCM90 at 0.025 C and (i) and (m) P-NCM90 at 0.5 C, (h) and (l) S-NCM90 at 0.025 C and (j) and (n) S-NCM90 at 0.5 C, (o) S-NCM70 at 0.5 C, and (p) S-NCM80 at 0.5 C.³⁴¹ Copyright 2021, American Chemical Society.

However, the structural characteristics of single-crystal Ni-rich cathodes also pose some problems, leading to controversial results in current single-crystal designs. The large size (∼µm) of the single-crystal particles leads to a long Li⁺ diffusion path and there is a lack of fast diffusion paths along the boundaries of the single-crystal particles, which results in slower Li⁺ diffusion kinetics and reduced rate performance.³⁴⁰ This will also lead to an inhomogeneous Li⁺ distribution, generating inhomogeneous structural stresses and inducing the formation of intragranular cracks. For example, Ryu et al.³⁴¹ systematically investigated the performance degradation and fading mechanisms of single-crystal NCM cathodes (S-NCM70, S-NCM80, S-NCM90) and poly-crystal NCM cathodes (P-NCM70, P-NCM80, P-NCM90) with the same particle size and Ni content of 70, 80, and 90%, respectively. They found that adopting a single-crystal cathode design exacerbated capacity fade and markedly impaired rate capability (Fig. 20d), while the Li diffusion coefficients did not differ significantly (Fig. 20e), suggesting that the decrease in rate performance is mainly due to the difference in the morphology of the single-crystal and poly-crystal cathodes. The lithium ions in single-crystal NCM cathode particles migrate by overall diffusion mainly through a two-dimensional path along the layer planes, whereas the poly-crystal NCM cathode has a three-dimensional fast diffusion path through the grain boundary network. In addition, they found that the capacity decay mechanism of the single-crystal NCM cathode is different from that of the poly-crystal NCM cathode. In situ XRD tests show that the single-crystal cathode experiences more severe structural inhomogeneity at a high rate due to long diffusion paths, with the presence of two-phase coexistence, and this phenomenon is exacerbated with increasing Ni content (Fig. 20g–p). It is further found that different regions of the charged S-NCM90 cathode particles have different lattice parameters, and the inhomogeneous structure induces inhomogeneous stress (Fig. 20f). This stress will lead to planar gliding and particle rupture of the cathode during the cycling process, which ultimately leads to capacity degradation. Therefore, the modification of the Ni-rich cathode by single-crystallization still needs to be assisted by other modification means to realize the comprehensive improvement of battery performance.

Another structural engineering strategy is core shell structure and concentration gradient design.³³⁴ On the surface of Ni-rich materials, high activity and high valence Ni³⁺/Ni⁴⁺ is easy to react with the electrolyte, resulting in the degradation of the surface structure of the material. Therefore, the design of a high to low concentration gradient of Ni ions from the core to the particle surface can effectively enhance the interface stability. The design principle is shown in Fig. 21a.³⁴² There are three generations of models (Fig. 21b–d). The first-generation concentration gradient design model uses NCM811 as the core, and the external shell is an NCM material with low Ni content and concentration gradient (Fig. 21b).³⁴³ The second-generation model is a full concentration gradient Ni-rich cathode material with LiNi_0.86Co_0.10Mn_0.04O₂ inside and LiNi_0.70Co_0.10Mn_0.20O₂ outside (Fig. 21c).³⁴⁴ This model solves the defect that the common core–shell structure is too thin to maintain stability at high temperature. The concentration gradient of Ni, Co and Mn is extended over the entire particle length. The third-generation model is a two-sloped full concentration gradient (Fig. 21d)³⁴⁵ LiNi_0.84Co_0.06Mn_0.09Al_0.01O₂ material with a double-oblique concentration gradient slope presenting different element content change rates from the inside to the outside, and the internal change is relatively gentle, while the external change is relatively fast. Recently, Kim et al.³⁴⁶ coated Co(OH)₂ on the surface of Ni_0.91Co_0.06Mn_0.03(OH)₂ (NCM91), and obtained a cathode material with rich Ni inside and poor Ni outside after sintering with lithium (Co-NCM91). In full cell tests, Co-NCM91 achieved a capacity retention rate of up to 85% after 300 cycles. Li et al.³⁴⁷ synthesized a NCM811 (CS) cathode material with a spinel shell by co-precipitation. The spinel shell has a three-dimensional Li⁺ transport channel, which can enhance electrochemical performance. The CS-003 sample with 3% spinel Li–Mn–O as the shell is charged and discharged at a high rate of 10 C, and the capacity retention is 81.6% after 200 cycles, which is much higher than the capacity retention of NCM811 which is 73.1%.

Fig. 21 (a) Design model of the Ni-rich cathode concentration gradient structure.³⁴² Copyright 2018, Elsevier. (b) First generation concentration gradient design.³⁴³ Copyright 2009, Springer Nature. (c) Second generation concentration gradient design.³⁴⁴ Copyright 2012, Springer Nature. (d) Third generation concentration gradient design.³⁴⁵ Copyright 2016, American Chemical Society. (e) Comparison of electrochemical properties of NCM811, LHEO and LHECNM.³⁴⁸ Copyright 2025, John Wiley and Sons. (f) Comparison of electrochemical properties of LNMO-800, CS-700, CS-750 and CS-800.³⁴⁹ Copyright 2023, Elsevier. (g) Comparison of cycling performance of W-containing core–shell materials with pristine and W-containing LiNiO₂.³⁵⁰ Copyright 2022, American Chemical Society.

Zhao et al.³⁴⁸ constructed a high-entropy gradient core–shell architecture (LHECNM) consisting of a Ni-rich, Co-free LiNi_0.89Mn_0.11O₂ core enveloped by a LiNi_1/6Mn_1/6Al_1/6Ti_1/6Mo_1/6Ta_1/6O₂ shell. The six-metal high-entropy shell exploits the cocktail effect of multiple cations to create a robust, concentration-gradient interface that passivates the surface against electrolyte attack and arrests TM migration. Compared with NCM811 and high entropy-doped Ni-rich cobalt-free LiNi_0.8Mn_0.12Al_0.02Ti_0.02Mo_0.02Ta_0.02O₂ cathode material (LHEO), this dual modification yields a discharge capacity of 201.6 mAh g⁻¹ at 0.3 C, retains ≥92% of its capacity after long-term cycling, and sustains stable high-rate performance up to 5 C, demonstrating the potency of combining compositional and gradient engineering (Fig. 21e). Jing et al.³⁴⁹ adopted a different but equally instructive route, first synthesizing a core–shell hydroxide precursor Ni(OH)₂@Mn(OH)₂ with a Ni-rich core and a Mn-rich shell. After lithiation at 700–800 °C, the Mn surface layer gradually diffused inward, producing a graded Li_1.08Ni_0.9Mn_0.1O₂ cathode whose gradient sharpness could be tuned by calcination temperature. At 700 °C (CS-700) and 750 °C (CS-750), a distinct Mn-rich outer layer persists, effectively buffering the high-Ni core from electrolyte contact; at 800 °C (CS-800), complete Mn/Ni homogenization erodes this protective gradient and accelerates degradation. Electrochemical data corroborate this trend: CS-750 delivers 187 mAh g⁻¹ at 0.1 C, and retains 91% after 100 cycles at 1 C and 79% after 500 cycles, whereas CS-800 suffers severe capacity fade linked to irreversible phase transitions near 4.2 V (Fig. 21f). The dQ/dV profiles further reveal that gradient-preserved samples (CS-700 and CS-750) maintain sharp, stable redox peaks, underscoring the critical role of a persistent Mn-rich gradient in suppressing structural degradation under high-voltage cycling.

In order to prevent elements such as Al, Mg, or Mn located in the shell from diffusing into the core during the heat treatment process of the core–shell structure, Rathore et al.³⁵⁰ used dry-coating of W to arrest inter-diffusion and stabilize thin shells. It was found that merely 1 wt% W, dry-coated onto hydroxide precursors, forms immobile WO_x clusters that physically block Mn (or Al) migration across core–shell boundaries during 800 °C calcination. In W-CS-NiMn-17/0.5 sample, this “pinning” effect preserves the intended Mn-rich outer shell; after 100 cycles, capacity retention is ∼80% higher than in the W-free analogue (Fig. 21g). The result demonstrates that sub-monolayer pinning agents can substitute for thicker, heavier coatings while still enforcing compositional gradients.

5. Summary and outlook

Degradation of the Ni-rich cathode structure (including the surface, interface and bulk phase) is mainly categorized into five distinct categories. (1) The formation of RLCs. This is due to the excessive lithium precursor introduced during synthesis or surface remodeling during cooling and storage stages. (2) Interfacial reactions. During cycling, the TM ion is constantly dissolved, and TM fluoride fragments form irreversibly. Surface RLCs also undergo side reactions with the electrolyte, producing organic and inorganic by-products that ultimately accelerate the growth of a CEI layer on Ni-rich particle surfaces. (3) Irreversible phase transition. The causes of irreversible phase transition on the surface of Ni-rich cathodes are complicated, which are related to cation mixing, oxygen release, charge/discharge rate, and SOC. (4) Microcrack generation. This is mainly caused by the anisotropic expansion/contraction of the Ni-rich cathode due to the irreversible phase transition, and is closely related to the SOC, pores and oxygen vacancies. (5) Cation mixing. The reasons for cation mixing include the similarity of Li/Ni ion radius, the low migration barrier of Ni, the effect of magnetic interaction and the poor thermal stability of Ni-rich cathodes.

The gas release of Ni-rich cathodes is mainly due to the decomposition of RLCs, degradation of the electrolyte on the surface, and the release of lattice oxygen under high voltage. As the voltage continues to rise, the surface RLCs formed during the preparation and storage process and the organic electrolyte will gradually decompose on the Ni-rich cathode surface to produce gas (e.g., O₂, CO and CO₂). In addition, at high SOC, the oxygen evolution potential is reduced, and the lattice oxygen is unstable and would eventually release in the form of O₂ gas. The above problems could be solved by different modification methods, including surface modification, element doping, electrolyte design, structural engineering, etc.

To further improve the energy density and stability of Ni-rich layered oxides and promote their large-scale application, future work should focus on the following aspects. Firstly, the selection of dopants and coatings should be assisted by further theoretical research. Secondly, it is necessary to achieve the controllable preparation of high-performance Ni-rich layered oxides with high lattice oxygen stability and high structural stability through structural engineering. Finally, it is important to implement composite strategies (particle engineering and surface/bulk stabilization techniques) to achieve synergies. Through continuous ingenious design and mechanism research, Ni-rich cathodes will have a wider application market in the future.

Conflicts of interest

There are no conflicts to declare.

Data availability

Data availability is not applicable to this article as no new data were created or analyzed in this study.

Acknowledgements

The authors acknowledge financial support from the National Natural Science Foundation of China (22408239) and the Fundamental Research Funds for the Central Universities. S. Jiang was supported by the Postdoctoral Innovation Talent Support Program (BX20240038).

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Footnote

† These authors contributed equally to this work.

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