Advances in doping strategies for sodium transition metal oxides cathodes: A review

Abstract

The electrochemistry of cathode materials for sodium-ion batteries differs significantly from lithium-ion batteries and offers distinct advantages. Overall, the progress of commercializing sodium-ion batteries is currently impeded by the inherent inefficiencies exhibited by these cathode materials, which include insufficient conductivity, slow kinetics, and substantial volume changes throughout the process of intercalation and deintercalation cycles. Consequently, numerous methodologies have been utilized to tackle these challenges, encompassing structural modulation, surface modification, and elemental doping. This paper aims to highlight fundamental principles and strategies for the development of sodium transition metal oxide cathodes. Specifically, it emphasizes the role of various elemental doping techniques in initiating anionic redox reactions, improving cathode stability, and enhancing the operational voltage of these cathodes, aiming to provide readers with novel perspectives on the design of sodium metal oxide cathodes through the doping approach, as well as address the current obstacles that can be overcome/alleviated through these dopant strategies.

Graphical abstract

Keywords

sodium-ion batteries / transition metal cathode / doping strategy

1 Introduction

The expenditure on global energy has experienced a significant growth in the past 20 years, hence creating the need for the advancement of energy storage systems that possess enhanced performance and sustainability. The efficient utilization of energy necessitates the implementation of energy storage systems that involve the conversion of energy from various forms, such as heat, radiation, electricity, chemistry, and gravitation, which are often challenging to harness.

Lithium-ion batteries (LIBs) have emerged as a highly promising technology for electrochemical energy storage in both mobile devices and utility scale applications since its initial commercialization by Sony Co. Ltd. in 1991 [1]. LIBs are embracing essential foundational elements, including Co, Li, Cu, and graphite, because of their widespread availability and significant economic importance [2,3]. It is expected that there will be material shortages and price hikes in the foreseeable future due to the reliance on LIB production on these crucial raw materials [4].

On the other hand, the electrochemistry of cathode materials for sodium-ion batteries (SIBs) differs significantly from LIBs and offers distinct advantages. Specifically, a multitude of well-defined cathode materials for SIBs have been synthesized through extensive investigations conducted by a diverse group of international experts [2]. The materials encompassed in this category consist of layered transition metal oxide and polyanionic compounds. Examples of ternary cathodes include NaTiO₂, NaVO₂, NaNiO₂, NaFeO₂, NaCrO₂ [5], along with mixed metal dioxides derived from them, such as NaNi_1/4Fe_1/4Co_1/4Mn_1/4O₂, NaNi_1/3Fe_1/3Mn_1/3O₂, NaFe_1/2Co_1/2O₂, and NaNi_1/3Fe_1/3Co_1/3O₄ [6–9].

Overall, the progress of commercializing SIBs is currently impeded by the inherent inefficiencies exhibited by these cathode materials, which include insufficient conductivity, slow kinetics, and substantial volume changes throughout the process of intercalation and deintercalation cycles. Consequently, numerous methodologies have been utilized to tackle these challenges, encompassing structural modulation, surface modification, and elemental doping.

This mini-review aims to highlight fundamental principles and strategies for the development of sodium transition metal oxide cathodes. Specifically, it emphasizes the role of various elemental doping techniques in initiating anionic redox reactions, improving cathode stability, and enhancing the operational voltage of these cathodes, aiming to provide readers with novel perspectives on the design of sodium metal oxide cathodes through the doping approach, as well as address the current obstacles that can be overcome/alleviated through these dopant strategies.

2 Fundamentals

2.1 Structures of layered transition oxide material-structures

Delmas’ classification scheme for layered alkali metal oxides designates sodium transition metal oxides as O3, P3, or P2, as per the established nomenclature as can be seen from Fig.1, in which the yellow atom is sodium which can be partially or fully occupied. The red denotes the oxygen atoms, and the green represents transition metal, in this case green, gray and purple represent different transition metals occupying the same site. Within this classification system, the letter denotes the specific alkali metal environment, with “O” representing octahedral and “P” representing prismatic. The numerical value, on the other hand, indicates the quantity of MO₂ slabs present in the hexagonal unit cell, which is 2 or 3 based on the notation.

The existence of ordered superstructures in alkali metal layers can be attributed to the strong interactions between Na⁺ ions [10]. The combination of this phenomenon, together with the arrangement of charges inside the layers of transition metals, has the potential to induce various phase transitions and voltage plateaus. The transport kinetics of Na⁺ are decelerated at the voltages corresponding to the phase shifts mentioned [11]. Moreover, these phase changes often exhibit incomplete reversibility, leading to a reduction in the overall cycle life. In contrast to a solid solution transition characterized by a sloping voltage profile, the absence of phase transitions in the system maintains the sodium transport kinetics at a relatively stable level, hence leading to prolonged cycle lifetimes [12].

Oxides with a sodium layered structure of the O3-type commonly experience many phase changes during the process of electrochemical sodiation and desodiation. These transitions include the transformations from the O3 phase to the O′3 phase, and further to the P3 phase and P′3 phase. Additionally, these oxides have sluggish insertion kinetics when highly sodiated [13]. As an example, when charged from 1.5 to 4.0 V, a O3 cathode undergoes a continuous O′3–P′3–P''3–O''3–O3 phase transition, with the majority of these transitions taking place at 3.0 V. The potential consequences of substantial lattice distortion and volumetric alterations in the crystal structure are thermodynamically adverse effects of multiphase evolution, which may compromise the overall structural integrity of the cathode. Additionally, electrolytes such as sodium hexafluorophosphate and fluoroethylene carbonate, have the capability to undergo disintegration and produce hydrogen fluoride at the interface between the electrode and electrolyte while the battery is in operation. The formed hydrofluoric acid attacks the transition metal (TM) ions with great ferocity, leading to the formation of TM fluorides and the subsequent irreversible loss of active mass from the cathode.

On the contrary, P2-type layered oxides exhibit a wide range of solid solution behavior and have rapid kinetics for the insertion and removal of sodium ions, and the P2 phase commonly shows the P2–O2 phase transition. The difference in the phase transitioning mechanism between P2 and O3 phase can be attributed to the extensive diffusion pathways provided by the prismatic sodium sites in the P2 structure [14].

2.2 Key challenges

As briefly discussed above, the biggest challenge is the irreversible transition of either the P2 or O3 cathodes: it is seen that both the P2 and O3 phases frequently undergo a succession of phase transitions characterized by varying arrangements of the oxide layers throughout the process of electrochemical cycling. The P2 phase often undergoes a transformation to the O2 phase as a result of the gliding (π/3 rotation) of particular metal oxygen octahedral polyhedrons when a specific quantity of Na⁺ ions are removed. This transformation leads to a notable constriction of the crystal structure and a reduction in the interlayer distance.

Another major challenge is the moisture sensitivity of the cathode in exposure to air. These cathodes can easily undergo reaction with water in the ambience, which results in possible water intercalation. There also could be {\mathrm{C}\mathrm{O}}_3^{2-} insertion, which could be subjected to further reaction. Sodium oxide or sodium carbonate species could also be generated from the surface, resulting in increased interfacial resistance during cell cycling. Section 3 will focus on the doping strategies to further overcome these particular challenges, and key examples will be used to illustrate the mechanism by which doping could alleviate the aforementioned issues.

3 Doping strategies

As mentioned before, both P2 and O3 structures have a challenging phase transition during the Na⁺ intercalation and de-intercalation process, which will harm the electrochemical performances to a large extent. With doping elements including Fe, Cu, Co, and F used, the Na⁺/vacancy ordered structure is eliminated and the interlayer distance is shortened. As a result, the detrimental phase transition is successfully controlled, the structure is mostly stabilized, and the air stability and electrochemical performances of the structure are significantly improved.

3.1 Single cation doping

3.1.1 Transition metal-ion substitution

Inactive elements: P2-type layered transition metal oxide, when compared to O3 structures, displays enhanced structural resilience due to the broader prismatic channels in TMO2 layers and the straightforward Na-ion diffusion. However, P2-layered transition metal oxide usually experience an unwanted P2–O2 phase shift when exposed to high voltages and having low sodium content, resulting in a substantial volumetric change and a quick depletion of capacity [15]. Moreover, the detrimental Na/vacancy ordering will bring multiple charge/discharge plateaus and harm electrochemical performance, which is more prevalent in sodium-ion cathode since Na has a larger ionic radius and will cause strong Na⁺–Na⁺ repulsion [16]. By introducing inactive elements with no d electrons (such as Li, Mg, Ti or Zn) into the TM layers, the metal lattice composition can be adjusted, Na/vacancy ordering can be suppressed, and Na diffusion can be promoted.

On the other hand, O3-type cathode material has a larger sodium content and a higher energy density, but it suffers from poor charge and cycle life due to the slow kinetic behavior of the narrow tetrahedron position and the reversible phase evolution of the migration of sodium ions. Moreover, the O3 phase frequently experiences a more complicated phase transition than the P2 structure, such as O3–O’3–P3–P3’ transition. P2-type Na_0.67[Ni_0.33Mn_0.67]O₂ cathode material is a promising choice for Na-ion batteries with high voltage Ni^4+/Ni²⁺ redox couple. However, it often undergoes unreversible and harmful P2–O2 phase transformation at a high voltage and causes significant capacity loss. Xu et al. [17] first evaluated the Li-substituted P2-type Na_x[Li_yNi_zMn_1–y–z]O₂ with nuclear magnetic resonance spectroscopy. Analyzing the local structure of this cathode material, they verified that Li-ions are located around Mn-ions in transition metal sites (Fig.2(a)). With Li-ions in the TM layer, more Na can stay in the prismatic site since it replaces mostly Ni sites and Li is at a lower valence state compared to Ni. Structural stability can also be maintained, thus delaying the inevitable P2–O2 phase transformation when all Na-ions are extracted from the P2 phase. This phase change will not occur even when charged to 4.4 V.

Similarly, Ghosh et al. [18] made an attempt using P2-Na_0.67+a[Li_xNi_0.33–yMn_0.67–z]O₂ cathodes. In this situation, the Ni⁴⁺/Ni²⁺ and Mn⁴⁺/Mn³⁺ redox couples both contributed during Na (de)intercalation. The ideal Li replacement reduces the Na⁺/vacancy ordering and suppresses the P2–O2 phase change. The improved P2-type cathode showed a good cycling stability and a large reversible capacity of 168 mAh/g at a low C rate.

Li substitution can also be effective in layered oxide cathode materials of the O3-type. Li-doped NaLi_0.1Ni_0.35Mn_0.55O₂ was synthesized by Zheng et al. [19] and characterized using particle X-ray diffraction (XRD) patterns. By refining the XRD patterns (Fig.2(b)), the evolution of the lattice parameters during the first charge/discharge process was determined. When the discharge process begins, the O3 phase and O′3 phase vanish rapidly, leaving just the P3 phase in the material. The P3 phase cannot transform into the O3 phase upon discharge, indicating that the O3 phase is thermodynamically stable.

Mg-ion doping is generally considered a candidate for transition-metal sites doping as well. Billaud et al. [20] found that Na_0.67Ni_1–xMg_xO₂ with Mg substitution can exhibit a superior structural stability by replacing the Mn³⁺ ions with Mg²⁺ ions, which have a high predilection for octahedral sites in the layered oxide framework. This can contribute to a highly stable structure with an initial capacity of up to 175 mAh/g.

Wang et al. [21] substituted some of the Ni-ions with Mg-ions for Na_0.67Mn_0.67Ni_0.33O₂ cathode materials and demonstrates that the nickel substituted cathode can significantly suppress the P2–O2 phase transformation. Similar to Li doping, the substitution of Mg²⁺ drives more Na⁺ ions to locate at prismatic sites and can stabilize the overall charge balance. Fig.2(e) illustrates the structural variations in Mg-substituted materials upon Na⁺ extraction and insertion and shows that the P2 phase can be maintained at a high voltage due to the larger number of Na in the prismatic site. Consequently, it displays a reversible capacity at over 123 mAh/g and an excellent capacity retention.

In addition, Siriwardena et al. [22] evaluated the effect of Mg doping on triclinic Na₂Mn₃O₇ transition metal oxide. The peak position at approximately 3.5 V in the cyclic voltammetry (CV) profile seems to smoothen when there is Mg-ion-doped, which shows that Na⁺/vacancy ordering is inhibited by Mg²⁺, which also enhances Na⁺ diffusion. After 30 cycles, the material with 2 mol.% Mg doping showed a high specific capacity of 143 mAh/g and a rate capability of 93 mAh/g. In addition, according to galvanostatic charge/discharge (GCD) analysis, O²⁻/n redox is extremely stable up to at least 90 cycles.o at least 90 cycles.

According to Feng et al. [23], Mg substitution can also optimize the electrochemical performance for P2-type NNMO cathode material. As mentioned before, Mg²⁺ which is situated in transition metal sites can induce more Na⁺ into the prismatic sites and can stabilize the charge balance. Since the low valence state Mg²⁺ can occupy nickel or manganese position, which is high in charge, the formed [Mg²⁺O₆] polyhedron can have more charges, leading to a stronger electrostatic attraction for Na-ions, thus stabilizing the structure (Fig.2(c)). The sample of Na_0.67Ni_0.18Mg_0.15Mn_0.67O₂ with Mg doping had an initial capacity of 123 mAh/g when discharged at a rate of 0.1 C between 2.0 and 4.3 V. After being used for 100 cycles at the same discharge rate, it retained 92% of its capacity.

When doping with Al, the interlayer spacing can be modified and the Jahan-Teller effect can be mitigated. Ramasamy et al. [24] introduced Al into P2-type Na_0.5Mn_0.5–xAl_xCo_0.5O_2. This in turn leads to a reduction in Mn³⁺ concentration, which minimizes the Jahn-Teller distortion, enlarges the distance between Na layers, and inhibits phase transitions. Similarly, the sodium layer spacing can also be enlarged by Al doping in O3-type NaNi_0.5Mn_0.5O₂ and the diffusion kinetics can be enhanced, according to Peng et al. [25]. Fig.2(d) represents the crystal structure diagram of O3-NNAMO viewed from [010] crystallographic axis and O3-NNMO itself. It is seen that the Na interlayer spacing is enlarged, which is beneficial for Na diffusion.

In most doping attempts, researchers are trying to strengthen the transition-metal layers by the strong bond between some inactive elements and oxygen and prevent rearrangement and sliding. For example, Park et al. [26] doped Ti into a P’2-type Mn-based cathode material due to the strong bond energy between Ti and O thus the stronger Mn-O-Ti-O-Fe–Mn-O-Ti-O bond in the transition metal layers may restrain the motions of Mn-O and Fe-O by sharing oxygen with Ti. According to the predicted crystal structure of Na_x[(Mn_0.78Fe_0.22)_0.9Ti_0.1]O₂ (Fig.2(f)), the length of c-axis will not change much even a large amount of Na are de/intercalated from/into the structure, which is hypothesized that the weaker Jahn-Teller distortion caused by Mn³⁺ is what causes the subtle structural change in Na_x[(Mn_0.78Fe_0.22)_0.9Ti_0.1]O₂ throughout the charge/discharge process.

As previously stated, Na/vacancy ordering can be lethal to electrochemical performance. Ti-doping can generate an entirely disordered Na-vacancy arrangement within Na layers, which will significantly alleviate this issue. Wang et al. [15] synthesized a Ti-doped P2-Na_2/3Ni_1/3Mn_1/3Ti_1/3O₂ cathode material that had an extended cycle life with 83.9% capacity retention after 500 cycles at 1 C and an exceptional rate capability. Simulations reveal the Na-ion routes in P2-Na_0.57NMT (Fig.2(g)), indicating that Na-ions travel along two-dimensional (2D) channels with associated diffusion pathways, resulting in rapid ion mobility. Achieving high Na⁺ mobility and low activation energy barriers results in an electrochemistry devoid of voltage plateaus.

Active elements: Although doping with those inactive elements can stabilize the structure, the specific capacity can be reduced due to their inert chemical property. Therefore, the substitution of active elements is introduced to attain the aims of inactive doping while maintaining the specific capacity.

Cu substitution can devote the Cu³⁺/Cu²⁺ redox couple which will enhance the working voltage. The cycle stability can also be improved with Cu doping. Kang et al. [27] investigated the Cu-doped P2-type Na_0.67Cu_xMn_1–xO₂ cathode material and the reaction potentials, capacities, and cycle performances are evaluated by CV. It can be observed from Fig.3(a) and 3(d), that Mg and Zn are not anticipated to engage in the electrochemical process, in contrast to Cu, which can undergo a Cu²⁺/Cu³⁺ reaction and exhibit a high potential plateau that can contribute to the rate performance. As a result, a capacity of more than 90 mAh/g is acquired at 12 C, and more than 70% of it can be obtained at this rate after 500 cycles.

Due to the similar ionic radii and valence states of Cu²⁺ and Ni²⁺, it was shown that the Cu substitution can not only improve the P2–O2 phase transition by encouraging more Na-ions to remain in the Na layers when charged to a high voltage but also improve air stability by preserving transition metal ordering while simultaneously reducing the capacity. Zheng et al. [28] conducted a study regarding Na_2/3Ni_1/3–xCu_xMn_2/3O₂ cathode material. It can be seen from Fig.3(b) that with Cu substitution, the voltage curve of the sample becomes smoother. The step-like voltage curve is eliminated which demonstrates that the sodium/vacancy ordering and phase transitions are suppressed to some extent. The schematic also shows that even when cycled to high voltage, the material still maintained the P2/OP4 structure, instead of harmful P2–O2 phase change. In the meantime, both Ni²⁺/Ni⁴⁺ and Cu²⁺/Cu³⁺ contribute to the redox reaction during cycling, contributing to less capacity during cycling. The potential reaction of Cu²⁺ with electrolytes is likely to be electrolyte-dependent.

Yang et al. [29] conducted a study regarding Fe-doped Na_2/3Ni_1/3Mn_7/12Fe_1/12O₂ cathode material to address the problem that NNMO often exhibits poor cycling stability triggered by the unwanted P2–O2 phase transition. According to Fig.3(e), the left side ex-situ XRD patterns for the Fe-doped sample detects the pure P2 phase, indicating that the P2–O2 phase transition is inhibited.

Researchers also doped Fe-ions into O3-type layered oxide cathode materials. Wang et al. [30] innovatively studied Na[Fe_1/3Ni_1/3Ti_1/3]O₂ using X-ray diffraction and absorption studies. It is found that besides the main redox couple Ni²⁺/Ni⁴⁺, the Fe²⁺/Fe³⁺ redox couple also contributes to the capacity. According to Fig.3(c), a sloping plateau in the charge/discharge profile and a minor peak in the dQ/dV plot can be seen at potentials below 2.0 V, which are related to the activity of Fe-ions, indicating that Fe- ions also participate in redox reactions.

Co substitution can also affect interlayer spacing, phase transformation, and rate capacity. Bucher et al. [31] first investigated the effect of Co-doping upon P2–Na_xCo_yMn_1–yO₂ using operando synchrotron XRD and in situ electrochemical impedance spectroscopy. The suppression of a Jahn-Teller-induced structural transition from the initial hexagonal to an orthorhombic phase is seen in Na_xMnO₂. At the same time, the suppression of Na⁺ ordering processes, and the enhancement of Na⁺ kinetics is also shown in measurements, which can lead to the higher cycling stability of the material.

Li et al. [32] examined the impact and functionality of Co substitution on the electrochemical performance and structure of Na_0.7Mn_0.7Ni_0.3xCo_xO₂. According to Fig.3(f), the replacement of Co³⁺ for Ni²⁺ increased the lattice parameter c and expanded the d-spacing of the Na-ion diffusion layer, which boosted the Na-ion diffusion coefficient and cathode material high-rate capacity. In addition, Co³⁺ substitution reduced the TM-O bond lengths and enhanced the structure stability.

Co-doped P2-type Na_0.67MnO₂ was presented by Fu et al. [33]. The schematic illustration described the XRD refinement data, indicating that the d-spacing is enlarged after Co doping, which will facilitate Na diffusion (Fig.3(g)). Moreover, the presence of Mn³⁺ will bring the Jahn-Teller effect and cause the irreversible distortion of the P2–P2’ phase change. With the Co substation, the portion of Mn⁴⁺ will increase and the Mn³⁺ level will lessen, thus alleviating phase distortion and the Jahn-Teller effect.

3.1.2 Alkali-metal substitution

In spite of modification at transition metal sites, alkali-metal site substitution has gained interest due to its ability to improve structural stability by enhancing electrostatic cohesion between adjacent TMO₂ layers and producing a pinning effect [13].

In 2014, Bae et al. [34] first disclosed the role of stabilizing the structure for Ca doping in O3-type NaNi_0.5Mn_0.5O₂. Due to the vast differences in ionic radius between Ca²⁺ and Ni²⁺ and Mn⁴⁺, Ca²⁺ ions are effectively incorporated into NaO₆. It is believed that the strong interaction of immobile Ca²⁺ with O₂ contributes to the restoration of the original O3 structure, resulting in the maintenance of distinct phase transition steps on discharge profiles after repeated cycles, which explains the excellent capacity retention of Ca²⁺ doped batteries.

Similarly, Matsui et al. [35] doped Ca²⁺ into P3-type Na_xNi_1/3Mn_1/3Co_1/3O₂ and surprisingly discovered that with Ca substitution, the irreversible O3′–O1 phase transition could be delayed at a high voltage. In addition, the material exhibits great cycling performance since it can suppress Na/vacancy ordering. A high-energy O3-Na_1−2xCa_x[Ni_0.5Mn_0.5]O₂ cathode was investigated by Yu et al. [36] and Ca²⁺ ions were also doped into NaO₆ octahedron. In contrast to Ca-free material, O3-type Na_0.98Ca_0.01[Ni_0.5Mn_0.5]O₂ displays a reversible O3–P3–O3 phase transition with little volume changes owing to the strong interaction between Ca²⁺ and O2.

Shen et al. [37] disclosed the unexpected effect of Ca²⁺ doping into sodium sites in the P2-Na_0.76Ca_0.05[Ni_0.23Mn_0.69]O₂ cathode material that can both initiate the anionic redox reaction and prevent phase distortion. Ca²⁺ tends to diffuse into the Na⁺ site and can function as a “pillar” to stabilize the structure, as depicted in Fig.4(a). The local electron distribution is generated to demonstrate that nonbonding oxygen 2p orbitals will be created from transition metal vacancies, and that these will significantly stimulate the anionic redox reaction. Due to the pinned Ca-ions in the Na sites, a robust layered structure with a suppressed P2–O2 phase transition is obtained.

Peng et al. [38] highlighted the peculiar alkali-metal layer modification with Zn²⁺ ions for P2-type Na_0.67Ni_0.35Mn_0.68O₂. The formation of an O2–Zn²⁺ –O2- “pillar” for enhancing electrostatic cohesion between two adjacent transition metal layers has halted the fragmentation of active material along the a-b plane and restricted the production of the O₂ phase. The crystal orbital Hamilton populations analyses demonstrated, as shown in lyses demonstrated, as shown in Fig.4(b), that Na-O had a shorter bond length and stronger bond strength than Zn-O, indicating an increase in the electrostatic cohesion of the alkali-metal layer. Fig.4(c) depicts the path of Na⁺ migration. In the Na layer, Na⁺ migrates across the 2D a-b planes, demonstrating the 2D diffusion paths. In contrast to pure Na_0.67Ni_0.33Mn_0.67O₂ (NNMO), the pathways of Na⁺ diffusion in Na_0.67Zn_0.05Ni_0.18Cu_0.1Mn_0.67O (NZNCMO) have a lower migration energy.

Instead of Ca²⁺ and Zn²⁺ substitution, K⁺ can also be doped into Na prismatic sites due to its large radius. Because of their shared electrical outer shells, potassium and sodium exhibit comparable chemical characteristics. Additionally, the bigger ion radius of the potassium ion than the sodium ions may raise the crystal lattice characteristics and impact the insertion/extraction process of the sodium ion. By doping potassium, the cycle ability and rate performance of multilayer cathodes in LIBs have been significantly enhanced [39].

Wang et al. [40] first synthesized K⁺-doped P2-type Na_0.67Mn_0.72Ni_0.14Co_0.14O₂. The PITT experiments show that the Na⁺ diffusion coefficients of K⁺ doped materials were increased throughout both the charge and discharge processes. Additionally, the Na_0.66K_0.01Mn_0.72Ni_0.14Co_0.14O₂ exhibited the greatest rate of Na⁺ diffusion. The lattice parameters changed after the substitution of K⁺.

Moreover, K⁺ ions were doped into P2-Na_0.612K_0.056MnO₂ in order to alleviate the gliding of transition-metal layers upon Na⁺ extraction and insertion by Wang et al. [41] According to Fig.4(d), the P2 phase is retained during the subsequent discharge procedure until 0.536 Na⁺ is injected for Na_0.65K_0.056MnO₂ (II). The P2 phase begins to arise with the coexistence of the P2 phase due to the small deformation of the MnO₆-octahedra slabs, and it becomes dominant in the fully discharged state (III) of Na_1.015K_0.056MnO₂.

The Mg²⁺ ion exhibits a significant disparity in ionic radius when compared to the Na⁺ ion, hence leading to its preferential occupation of transition metal sites during the process of doping. Nevertheless, under specific circumstances where a certain number of TM sites have already been occupied, Mg²⁺ ions can still occupy the alkali-metal site. In their study, Wang et al. [42] successfully synthesized P2-type Na_0.7[Mn_0.6Ni_0.4−xMg_x]O₂ materials by introducing Na site doping. It is evident from the observation that the introduction of doped Mg ions into the Na sites effectively inhibits the phase transition from P2 to O2, hence providing stability to the structure, especially in the extensively desodiated state. Magnesium ions can function as supportive “pillars” that effectively mitigate structural collapse in a specific direction under the influence of a high-voltage charge. On the contrary, the presence of Mg-ions in the sodium layer can lead to the formation of “Na-O-Mg” and “Mg-O-Mg” configurations. This arrangement results in the incorporation of ionic O 2p character, which causes these O 2p states to overlap with those that interact with transition metals in the O-valence band. Consequently, this interaction facilitates reversible oxygen redox processes.

Li et al. [43] were subsequently motivated to inject Mg into the Na layer of the P2-layered Na_0.7Mg_0.2[Fe_0.2Mn_0.60]O₂ cathode material. Within the TM layers, TM vacancies are effectively created to compensate for the charge difference. The crystal lattice configurations for a series of doped Na cathodes described are depicted in Fig.4(f). Mg²⁺ ions are effectively doped into Na sites, and the unit cell decreases continuously along the c-axis. Huang et al. [44] also considered the Mg²⁺ sodium layer substation for P2-Na_5/6Li_1/4Mn₄O₄. Similarly, Mg²⁺ in this situation can also function as a “pillar” and provide nonbonding O 2p orbital to activate anionic redox reaction. In addition, the expansion of Na_0.773Mg_0.03Li_0.25Mn_0.75O₂ during sodium insertion/extraction is only 1.1%, indicating a “zero-strain” cathode.

3.2 Single anion doping

In contrast to cation doping, there is insufficient research regarding anion doping for sodium-layered transition metal oxide cathode material. Recently nonmetal ions such as F^– and B^– are investigated to be doped in oxygen sites and can surprisingly give a superior electrochemical performance due to their strong electronegativity and less negative valence.

3.2.1 Non-metal anions

In 2017, Zhang et al. [45] first fabricated the F-doped O3-NaNi_1/3Fe_1/3Mn_1/3O₂ cathode material. It was verified that F-doping could modify the binding energy of oxygen and increase the rate of Na⁺ diffusion. Moreover, the cell characteristics initially declined and subsequently rose as a function of F doping, which was attributed to an increased fraction of Mn³⁺ rather than Mn⁴⁺ when O²⁻ was substituted by F⁻. The binding energy of TM-O boosted by F doping could aid in the suppression of the John-Teller effect of Mn³⁺ and the facilitation of Na⁺ diffusion. As a result, the NFM-F_0.01 material achieves an optimal cycling performance with a capacity of 110 mAh/g at a current density of 150 mA/g after 70 cycles.

Shi et al. [46] created the F-substituted Na_0.44MnO_1.93F_0.07 with a layer-tunnel hybrid structure using the same principle. Fig.5(a) demonstrates that F-ions are substituted into the O2-layer and convert a portion of Mn⁴⁺ to Mn³⁺ due to the decreased negative charge of oxygen. In addition, its reversible capacity is vastly superior to that of the tunnel phase Na_0.44MnO₂, with a discharge capacity of 149 mAh/g at 0.5 C and 138 mAh/g at 1 C. In addition, the electrode has exceptional cycling stability, retaining approximately 79% of its capacity after 400 cycles at 5 C.

On the other hand, Chen et al. [47] discovered that anion doping could boost the activity of Ni²⁺ by using F-substituted Na_0.6Mn_0.95Ni_0.05O_2–xF_x cathode. The charge/discharge profile revealed that the strong Ni−F bond due to F doping could maximize the use of Ni²⁺/Ni³⁺ redox couple since there is an obvious and high-intensity reduction peak around 3.1 V, indicating a highly reactive reduction reaction of Ni³⁺/Ni²⁺. Furthermore, the strong Na−F bond can prevent the collapse of the P2 phase structure and increase cycle stability. According to Fig.5(b), the Na_0.6Mn_0.95Ni_0.05O_1.95F_0.05 sample shows stable cycling and capacity.

A series of F-substituted Na_2/3Ni_1/3Mn_2/3O_2−xF_x cathode materials were synthesized and investigated by Liu et al. [48]. It is illustrated that after doping with F⁻, both Ni and Mn would activate in the redox reaction, which could bring an extra specific capacity of 10 mAh/g. Furthermore, F substitution could disturb the John-Teller effect during the cycle by decreasing the ordering of Ni/Mn. As seen in Fig.5(c), Na_2/3Ni_1/3Mn_2/3O_1.95F_0.05 exhibits the greatest rate performance with high reversible specific capacities. These exceptional rate capabilities are primarily due to better interface stability and, as a result, decreased resistance.

Similarly, Kang et al. [49] doped F- into P2-Na_0.6Mn_0.7Ni_0.3O_2–xF_x with a solid-state reaction. It is shown that both the structure and the electrochemical activity P2-Na_0.6Mn_0.7Ni_0.3O_2–xF_x can be altered through F doping. Meanwhile, the proportion of Mn³⁺/Mn⁴⁺ boosts from 0.72 to 1.61 for the increased amount of F doping to balance the charge compensation, according to the X-ray photoelectron spectroscopy (XPS) pattern of Mn 2p in Fig.5(d). F^– substitution can activate Mn⁴⁺ reduction and the increased amount of Mn³⁺ will occupy Ni²⁺ sites, disturbing the Ni²⁺/Mn⁴⁺ cation ordering and can decrease the Jahn-Teller effects thus further improving structural stability.

Recently, Liu et al. [50] conducted a study with F-doped O3-NaNi_1/3Fe_1/3Mn_1/3O₂. They stressed the unique characteristics of F^– regarding lightweight and strong electronegativity and pointed out that these properties could enhance the TM-O bond strength and widen the Na⁺ diffusion channel. In addition, the strong electronegativity effectively suppresses the slip of the interlayer and could enhance structural stability to a large extent. Fig.5(e) shows the nudged elastic band (NEB) calculation of activation barrier energy for Na⁺ diffusion in two structures. It is determined that after doping, the migration energy barrier is decreased by approximately 0.33 eV, which considerably enhances the migration rate of Na⁺ and is favorable.

F^– doping can also be adopted to O3-type Na[Ni_0.5Mn_0.5]O₂. Yu & Sun [51] optimized it with F^– substituted Na_0.95[Ni_0.5Mn_0.5]O_1.95F_0.05 and achieved a high capacity retention. Most importantly, they found that fluorination not only reduces the microcracking of cathode secondary particles but also prevents surface deterioration. As can be seen in Fig.5(f), the NM55 cathode undergoes a phase transition that induces stress when charging, leading to the formation of microcracks. These cracks promote electrolyte infiltration, enabling undesired surface reactions on primary cathode particles. Moreover, the edges of cathode particles show tearing, likely due to the dissolution of transition metals from oxygen release reactions, causing battery performance degradation. Fluorine substitution in the crystal structure of the cathode significantly mitigates this degradation. An appropriate amount of fluorination of the NM55 cathode prevents severe phase transitions, reduces microcracking, and enhances chemical stability. This results in an improved resistance to surface degradation and a reduced dissolution of transition metals. Wang et al. [52] found that B doping could convert the NaMNO from the P3 phase to a more stable P2 phase when they investigated B-doped Na_0.65Mn_0.75Ni_0.25B_0.1O₂ and it exhibited a better cycling life compared with that of the F-doped one.

Guo et al. [53] also introduced B-ions into O'3-type NaMnO₂ to occupy Mn sites. According to Fig.5(g), B³⁺ ions are accommodated in octahedral sites in the transition-metal layer. Na⁺ ions are found in the octahedral position, whereas Mn³⁺ and B³⁺ ions are found in the transition-metal layer. After B doping, the average effective radius of the transition-metal layer is lowered, making the structure more stable. Simultaneously, B prevents the Jahn-Teller distortion generated by Mn³⁺.

Moreover, Wang et al. [54] innovatively constructed a BO3 triangle configuration (Na_0.67Ni_0.3Co_0.1Mn_0.6O_1.94(BO3)_0.02) to alter the local structure of the P2-type material. This design acts as a strong support for the whole structure by stopping water insertion during Na (de)intercalation, thereby preventing deformation. Additionally, the arrangement allows more Na-ions to position themselves on the edge-sharing octahedrons, promoting better Na⁺ movement. As shown in Fig.5(h), following B-doping, additional Na occupies the Nae site, affecting the ratio of Nae/Naf and disturbing the intrinsic Na⁺/vacancy ordering, which will permit a smooth diffusion process to benefit the Na⁺ transit. As a result, the material showcases minimal volume change and an impressive capacity retention of 80.1% even after 1000 cycles at 2 C.

3.3 Multi-ion co-substitution

As the field of cathode development for SIBs progresses, multi-ion co-substituted layered oxides are emerging as an important avenue of research. Similar to single-ion doping, researchers have explored the incorporation of various cationic and anionic elements into cathodes. This diverse range of elements, including but not limited to Fe, Mn, Li, Co, Cu, Mg, Ti, and Ni [55–62], is introduced to harness their synergistic effects to enhance the electrochemical performance of the cathode, which is often a strategy employed in Li-ion cathode materials. The co-doping strategy, involving different combinations of these elements, has proven effective in bolstering the cyclic stability of the cathode, triggering reversible anionic redox reactions, and thereby elevating the overall electrochemical performance of SIBs batteries.

3.3.1 Preventing P2–O2 phase transition for P2-type cathode

Layered transitional metal-based oxide cathode with a P2 and O3 crystallographic structure is a promising class of cathode for SIBs. P2-type cathodes, in comparison to O3-type cathodes, shows a higher diffusion coefficient owing to open prismatic paths, high-rate kinetics, and high ion conductivity [63]. However, there are still defects for P2-type cathode due to deficient Na⁺ [64]. For Mn based P2-type cathode specifically, the interlayer-gliding MeO₆ can easily occur when Na⁺ is extracted at a deep state of charge. This results in the unwanted P2–O3 phase transition. This transition is not favored because of the large crystal volume expansion that leads to the degradation of crystal structure and capacity reversibility [63]. To address the aforementioned issues, the synergetic effect of multiple ions emerges as a practical solution.

Previous research has established that ion dopants such as Li-ions [17], Mg-ions [65], Zn-ions [65], Cu-ions [66], Ti-ions [67], and others [68,69] can inhibit the interlayer gliding of MeO₆ octahedral sheets. In a recent study conducted by Li et al. [63], an innovative approach has illuminated the promising effects of co-doping in manganese-based layered oxide cathode materials, as depicted in Fig.6(a). This strategy aims to suppress the P2–O2 phase transition by constraining the octahedral sheets and eliminating the Na⁺/vacancy ordering resulting from Na⁺–Na⁺ repulsion, thus reducing the dimensionality of sodium migration and the Na⁺ ion diffusion coefficient. This is achieved by employing a two-step doping strategy, where Co (Fig.6(e)) [70] is doped during the synthesis of precursors using a coprecipitation method to eliminate Na⁺/vacancy ordering, while Li doping (Fig.6(f)) [17] occurs during a solid-state reaction to mainly suppress P2–O2 transition. As anticipated, the Co/Li co-doping successfully eliminates the P2–O2 phase transition. Furthermore, due to the suppression of the P2–O2 transition, oxygen redox reactions, which were originally expected to occur in the O2 phase, now take place in the P2 phase. This results in different redox potentials and triggers reversible anionic red ox reactions at voltages exceeding 4.0 V (Fig.6(c)) [63], thereby reducing oxygen loss and preventing rapid capacity decay observed in undoped cathodes. Notably, the plateau observed in the curve of the undoped cathode in Fig.6(d) [63] below 4.0 V, attributed to Na⁺/vacancy ordering, disappears in the Co/Li co-doped cathode curve, indicating the elimination of Na⁺/vacancy ordering. Fig.6(b) [63] illustrates that this cathode exhibits a higher specific capacity compared to most other cathodes doped with various strategies [63].

Anilkumar et al. [64] have introduced an innovative approach involving Cu and Li co-doping to the P2–O2 phase transition and eliminate Na⁺/vacancy ordering. Building upon previous studies that demonstrated the effectiveness of Li⁺ doping in increasing Na⁺ content in P2-type layered cathodes to inhibit Na deficiency and suppress Na⁺/vacancy ordering, Li⁺ was selected as one of the dopants for this strategy. However, as the advantage conferred by Li-ion doping gradually diminishes and capacity degrades due to the loss of Li ions during cycling [17,72–74], Anilkumar et al. [64] opted to co-dope the layered oxide cathode with Cu²⁺ to stabilize the structure while maintaining a high Na content, ultimately achieving a high rechargeable capacity under high operating voltage. Adhering to the guideline of maintaining a cutoff ratio between the interlayer d-spacing of the transition metal and sodium layer at approximately 1.62 [75] (otherwise it transitions into the O3 phase), Anilkumar et al. [64] determined that the material doped with 0.07% Cu exhibited the best performance. As evident in Fig.7(b) [64], nearly all redox peaks align with each other, forming a solid phase with high reversibility. This solid phase effectively suppresses the P2–O2 phase transition, leading to high reversibility and cycling stability during Na⁺ intercalation/deintercalation. Other remarkable properties, such as high capacity retention, are demonstrated in Fig.7(a) [64]. This figure illustrates that the Cu–Li co-doped material maintains a capacity retention of approximately 75% (~85 mAh/g) after 200 cycles and a Coulombic efficiency of 99.6%. Furthermore, Fig.7(c) [64] showcases the impressive stability of this Cu–Li co-doped material, maintaining a highly reversible discharge capacity of ~110 mAh/g after 10 cycles.

Zhao et al. [76] have made significant advancements in the development of a novel P2-type manganese-based cathode by co-substituting Cu for Mn and Co for Ni, as illustrated in Fig.7(d) [72]. Active Co substitution efficiently suppresses the P2–O2 phase transition, while Cu doping [77], as indicated by previous research, enhances the reversibility of anionic redox reactions [67]. Furthermore, this construction strategy optimizes the lattice structure of TMO2 sheets, resulting in the shrinkage between TMO2 layer and Na layer [71]. This optimization inhibits undesirable P2–O2 phase transitions and reduces the barrier to Na⁺ migration. Co substitution elevates the high-voltage region to activate oxygen redox and weaken irreversible oxygen loss, while Cu substitution introduces interactions between Cu 3d and O 2p orbitals to suppress voltage decay. As the promotion of reversible redox reactions and the suppression of irreversible oxygen loss are crucial aspects of their strategy, Fig.7(e) [76] provides insights into the redox couples spanning the entire voltage range. The Cu/Co co-doped cathode demonstrates highly reversible anionic redox reactions within the high-voltage range of 4.2 to 4.5 V, significantly reducing polarization of the redox pairs and thereby enhancing electrochemical activity while mitigating irreversible oxygen loss. The voltage plot in Fig.7(f) [76] provides further evidence of the benefits of the Cu–Co co-doped cathode, as it demonstrates a higher initial Coulombic efficiency and reduced irreversible capacity loss.

3.3.2 Improving O3-type layered cathode performance

In contrast, layered oxide cathodes of the O3-type have distinct advantages over layered cathodes of the P2-type. O3-type cathode is a prospective cathode for SIBs due to its superior chemical activity, simple synthesis method, and greater charge/discharge capacity than P2-type cathode [77, 78]. However, when exposed to air, the O3-type cathode suffers from rapid capacity decay due to spontaneous extraction of Na, oxidation of transition metals, and concurrent transition to Na-deficient phases [79–83]. This poor air stability limits the widespread application of O3-type cathode as it raises the price of material storage, transportation, and battery production [83]. Moreover, because the O3-type cathode endures a reaction process involving numerous phase changes and volume variations, and because the O3-type cathode is additionally sensitive to moisture, its lifespan is significantly diminished [84]. To optimize the use of O3-type cathodes, researchers have exerted significant effort to surmount these obstacles.

Yao et al. [85] developed a strategy for the O3-type NaNi_0.5Mn_0.5O₂ cathode. This strategy involves modifying the structure of the cathode by reducing the distance between Na layers while simultaneously increasing the valence state of transitional metals. They selected a combination of Cu and Ti to achieve this modification, as shown in Fig.8(a) [85]. The choice of Cu and Ti was based on their ability to suppress spontaneous Na extraction and enhance the resistance of the material to oxidation. Ti was chosen as a co-dopant along with Cu because it could effectively increase the valence state of Ni through electronic delocalization, which reduced the number of electrons around Ni [86]. Cu was selected due to its electronegativity, electrochemical activity, and substantial difference in Fermi level compared to Ti [87]. As depicted in Fig.8(b) [85], this co-substitution strategy effectively inhibits undesired phase transitions, as confirmed by XRD patterns that show no change in morphology whether the material is placed in water or air for 2 h. Moreover, the improved air stability is evident in the significantly elevated capacity retention after aging tests. Most notably, this change results in significant improvement in air stability compared to the undoped cathode, while maintaining the original structure and capacity.

Employing the same elements as dopants, Wang et al. [84] sought distinct enhancements for O3-type oxide cathodes through a Cu²⁺/Ti⁴⁺ co-doping strategy. Previous research has established the advantages of Ti⁴⁺ doping in both P2 and O3 structures, particularly in augmenting rate capacity and bolstering cycling stability. Crucially, the heightened ionicity of the lattice structure, induced by Ti⁴⁺, contributes to the reduction of phase transitions, thereby enhancing stability across a broader voltage range [88–90]. Simultaneously, Cu substitution aids in improving the moisture stability of the cathode [85]. Consequently, Wang and colleagues [84] proposed the co-substitution of Cu and Ti in O3-type oxide cathodes. They synthesized two sets of Cu/Ti co-doped materials, with varying amounts of Ni being replaced by Cu: NaNi_0.5−yCu_yMn_0.4Ti_0.1O₂ (z = 0.1) and NaNi_0.5−yCu_yMn_0.3Ti_0.2O₂ (z = 0.2). Fig.8(c) presents the electrical evaluation, illustrating a progressively smoother curve as the Cu content increases, indicative of a diminishing phase transition with higher Cu content [84]. Furthermore, Fig.8(d) compares the retention performance between the co-doped O3 oxide cathode and the NVPF cathode, another exemplary SIB cathode [84]. Remarkably, both sets of cathodes not only surpass the undoped material but also outperform the NVPF cathode, promising enhanced cyclic stability.

Wang et al. [84] successfully developed a Ti/Zr-based O3-type cathode through solid-phase synthesis, i.e., NaNi_0.5–xMn_0.3TiZr_xO₂ (x = 0.02, 0.05). Zr is utilized to enhance cycling and rate performance due to its larger ion radius and stronger bond energy with O than NiO or MnO. Zr has a larger ion radius and a stronger bond energy with O compared to NiO and MnO [91]. This partial substitution of electrochemical Ti/Zr brings higher electronic delocalization and entropy of mixing that results in improved structural stability [92]. In addition, the increased interlayer spacing and excellent conductivity of the co-doped cathode contribute to the enhanced rate capacity. Fig.8(e) depicts the CV test outcomes of NaNMTZ2 at various scan frequencies, which differ little from the outcomes of NaNMTZ2 [92]. Co-substitution of Ti/Zr can inhibit multiphase transformation by inhibiting the gliding of transition metal oxide during Na intercalation/extraction. In addition, Fig.8(f) illustrates the cycling performance of the electrodes at 0.05 C and voltages between 2 and 4 V [92]. NaNMTZ2 has a capacity of approximately 62%, which still represents an enhancement over undoped material (57% retention).

4 Conclusions and perspectives

While the current literature has many strategies to solve existing issues associated with using sodium cathodes in a full SIB, this paper focuses on using dopants to solve existing issues associated the sodium cathodes. Specifically, single-ion doping using cation and anion, as well as strategies and cathode material using multi-ion doping is covered. Fig.9 displays a schematic representation of the structures and doping effects of sodium transition metal cathodes.

Specifically, the use of cationic doping is highly effective in preventing irreversible phase transitions, leading to significant improvements in the electrochemical performance of layered oxides [86–88]. The Jahn-Teller distortion can be mitigated in part by doping, while the interaction between transition metal and oxygen can be enhanced in part by facilitating oxygen sharing with stable metallic elements. Consequently, the motion of the TM-O6 octahedron is impeded, resulting in the suppression of the phase transition. Ca-doped P3-type Na_xNi_1/3Mn_1/3Co_1/3O₂ material effectively inhibits the phase transition from O′3 to O1, especially at high voltages greater than 4.0 V. Ca²⁺ and transition metal ions repel one another, resulting in an increase in the interlayer spacing within the sodium layer. This expansion increases the diffusion rate of sodium ions, thereby substantially improving the electrochemical performance of sodium during cycling [89].

Doping with various anions, such as F, Cl, and S_x, has been proposed. The participation of the anion facilitated partial charge neutralization in the overall redox reaction of the entire cell, and this process demonstrated both rapid kinetics and long-term stability. This proof of concept presents enticing opportunities for achieving greater reversibility. Doping has been observed to substantially increase the ionic conductivity of the material, reduce cation mixing, minimize oxygen release, prevent irreversible oxygen loss, stabilize the structure of the material, and improve its cycle performance. Therefore, the material demonstrated significant improvements in both rate performance and cycle performance. In addition, it has been discovered that F-doping is effective at regulating the binding energy of oxygen and the Mn^3+/Mn⁴⁺ ratio, thereby mitigating the Jahn-Teller distortion of Mn³⁺. In addition, fluorine doping has been found to increase the diffusion rate of sodium ions and improve the overall rate performance of the material. Notably, it seems that fluorine anion doping is particularly helpful in inhibiting irreversible oxygen loss and structural distortion. It could be due to the small size of fluorine and its highly electronegative character. Fluorine doping has the potential to induce surface reconstruction and alter the electronic structure through modifications in the covalence of metal-anion bonds and the displacement of the p-band center [90,91].

Multi-ion/co-doping strategies demonstrate that they can potentially address issues of the suppressing Jahn-Teller-induced structural transition and suppression of the ordering process of Na-ions. Co-doping has distinct effects from those intended for other transition metals. Utilizing the high potential of their respective redox reactions, the incorporation of elements such as Ni or Fe into the structure seeks to increase energy density. Moreover, the utilization of high entropy dual or multi-dopant cathodes allows for the possibility of ions of diverse sizes, which can enhance the kinetics of sodium diffusion and stabilize the structures. However, more efforts should be devoted to understanding the effects of multi-ion doping and their potential versus single-anion or cation doping. It remains crucial to elucidate its importance in terms of the optimal dopant concentration for the specific active site inside the host material. Moreover, it is necessary to verify the synergistic impact of multielement doping. The rational design of adapting appropriate concentrations of different ions should be elucidated. The recently popular high entropy concepts used in solid ionic conductors and Li cathodes which increase the ion diffusion can also be potentially used in designing these sodium cathodes.

Furthermore, the present investigation primarily utilizes metallic sodium as the counter electrode to assess the efficacy of the electrode within a half-cell. It is widely acknowledged that the performance of batteries in half-cells differs significantly from that of full cells featuring two insertion materials. The safety handling and the associated fabrication cost of Na metal should also be taken into consideration. Ongoing efforts are also required to optimize binders, electrolyte additives, and battery assembly technology to advance the practical application of Na full cells. In solid-state batteries, the interfacial formation with sodium metal and solid electrolytes is a direction for future study, as well as the mechanical properties of sodium metal and how it affects performance in a full cell. In general, the superior rate capabilities of Na metal anodes in comparison to Li metal anodes thus far may be attributed to the lower elastic moduli and plastic properties exhibited by Na in comparison to its lower melting temperature. In conclusion, it was discovered that interfacial forces between sodium metal and other solid-state battery components inhibit sodium deformation.

On the full cell level, SIBs still have major challenges such as their respective interfaces with existing liquid electrolytes and or solid-state electrolytes. Particularly, the interfaces between most of the sodium cathodes and the existing solid-state electrolytes are not well understood. A new electrolyte system with specific formulations, an optimized solvent, sodium compounds, and additives requires further investigation. As large-scale energy applications demand that SIBs perform well under more constrained climatic and intermittent conditions, the organic liquid electrolytes used in SIBs should have a wider temperature range tolerance while allowing for safe and stable cycling.

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