Structural transformation and electrochemical properties of a nanosized flower-like R-MnO₂ cathode in a sodium battery

Abstract

High-energy density and low-cost sodium–ion batteries are being sought to meet increasing energy demand. Here, R-MnO₂ is chosen as a cathode material of sodium–ion batteries owing to its low cost and high energy density. The structural transformation from the tunnel R-MnO₂ to the layered NaMnO₂ and electrochemical properties during the charge/discharge are investigated at the atomic level by combining XRD and related electrochemical experiments. Na_≤0.04MnO₂ has a tunnel R-MnO₂ phase structure, Na_≥0.42MnO₂ has a layered NaMnO₂ phase structure, and Na_0.04−0.42MnO₂ is their mixed phase. Mn³⁺ 3d⁴[t_2gβ3d_z²(1)3d_x²−y²(0)] in NaMnO₂ loses one 3d_z² electron and the redox couple Mn³⁺/Mn⁴⁺ delivers 206 mA h g⁻¹ during the initial charge. The case that the Fermi energy level difference between R-MnO₂ and NaMnO₂ is lower than that between the layered Na_(12-x)/12MnO₂ and NaMnO₂ makes the potential plateau of R-MnO₂ turning into NaMnO₂ lower than that of the layered Na_(12−x)/12MnO₂ to NaMnO₂. This can be confirmed by our experiment from the 1st–2nd voltage capacity profile of R-MnO₂ in EC/PC (ethylene carbonate/propylene carbonate) electrolyte. The study would give a new view of the production of sustainable sodium battery cathode materials.

---

1. Introduction

Sodium–ion batteries (SIBs) have been studied as low-cost (e.g., the cost of Na₂CO₃ is about 25–30 times cheaper than Li₂CO₃)¹ and material-abundant alternatives to lithium–ion batteries (LIBs) due to the abundant resource of sodium on earth² and similarities of intercalation chemistry between SIBs and LIBs.³ However, the larger ionic radius of Na (1.02 Å) than Li (0.76 Å) limits the choice of electrode materials.⁴ Cathode materials of SIBs such as organic compounds, Prussian blue, polyanionic compounds,^5–7 and oxides^8–10 have been well investigated. Among all the cathode materials, sodium layered oxides (Na_xMO₂), in which the structure can be easily regarded as MO₂ layers of edge-sharing MO₆ octahedron and Na⁺ occupying interlayer space, have attracted considerable attention.¹¹ Their theoretical capacity is 230–250 mA h g⁻¹ and depends on the M constituents.¹² Up to now, there has been a significant level of interest in sodium and manganese based layered oxides, due to their many advantageous properties, such as environmental friendliness and relatively less cost.¹ But, sodium layered oxide cathode materials still have some problems such as low average voltage, long-term cycling instability from the effects of Mn³⁺ Jahn–Teller distortion and multiple phase transitions, and poor rate capability.

A range of different Na_xMnO₂ (x ≤ 1) structures, from the tunnel to the layered, were synthesized and described by Hagenmuller's research group in 1971.¹³ Subsequently, a series of Na_xMnO₂ based materials, Mg-doping compounds (e.g., P2-Na_2/3[Mn_1−xMg_x]O₂(0 ≤ x ≤ 0.2)),¹⁴ Li-doping compounds (e.g., P2-Na_5/6[Mn_0.75Li_0.25]O₂)¹⁵ and binary Na_xMnT_MO₂ (T_M = Fe, Ni, Co and Cu)^16–18 and ternary Na_xMnO₂ ( = Fe, = Ni; = Fe, = Ti; = Ni, = Co; = Ni, = Cu; and = Ni, = Ti;)^19–23 systems and other Mn-based materials for LIBs have been studied.^24,25 In these materials Mn³⁺ 3d⁴ would give rise to the Jahn–Teller distortion of the Mn³⁺O₆ octahedra, generally leading to structural degradation on cycling.¹ For example, MnO₂ with α-MnO₂, β-MnO₂, γ-MnO₂, δ-MnO₂, λ-MnO₂ crystalline phases and so on as the cathode materials of SIBs^26–29 suffer the structural instability and Mn²⁺ dissolution (there is a reaction of Mn³⁺ → Mn²⁺ + Mn⁴⁺) in electrolyte during the cycling.^30–35

Theoretically, MnO₂ structural change with various tunnels during the Li⁺/Na⁺ ion insertion has been paid attention to. Tompsett and Islam³⁰ found that Li⁺/Na⁺ ions distributed themselves as uniformly as possible at the 8h or 8h′ sites of the 2 × 2 tunnel for α-MnO₂ and some Mn–O bonds are elongated to change MO₆ octahedrons to MO₅ tetragonal pyramids. Li et al.³⁶ found that the transformation energy barriers between the layered MnO₂ structure and MnO₂ with various tunnels are small. However, a distinct picture of the structural transformation from the tunnel to layered MnO₂ structures with the ion insertion content is absent, which is important to understand the electrochemical properties of MnO₂ during the charge/discharge.

Recently, we prepared nanosized orthorhombic R-MnO₂ with the 1 × 2 tunnel structure experimentally and studied its electrochemical properties.³⁷ R-MnO₂ is one among various crystalline phases besides α-MnO₂, β-MnO₂, γ-MnO₂, δ-MnO₂, and λ-MnO₂.³⁷ R-MnO₂ transforms from the tunnel to the layered NaMnO₂ after initial discharge, and the next discharge/charge is the transformation between the layered Na_0.33MnO₂ and the layered NaMnO₂.³⁷ The initial activation process and the structural change during the charging/discharging are similar to other Mn-based materials, such as Li[Li_0.2Ni_0.16Mn_0.56Co_0.08]O₂.³⁸ In this work, DFT calculations are carried out to thermodynamically study the structural transformation of MnO₂ from the 1 × 2 tunnel to the layered with the Na⁺ ion insertion content, and XRD experiments are combined to confirm the calculated result. To understand the electrochemical properties, the density of state (DOS), local DOS, and Bader charges of layered NaMnO₂ during the charge are investigated at the atomic level. Different potential plateaus of the 1st and 2nd voltage capacity profile in the EC/PC electrolyte from the experiment are explained by exploring Fermi energy level and electron density distribution of the highest occupied molecular orbital (HOMO) of R-MnO₂ and the layered Na_(12−x)/12MnO₂.

2. Calculation method and model

R-MnO₂ with the 1 × 2 tunnel structure is orthorhombic (Space group: Pnma), and optimized primitive lattice parameters are a = 9.51 Å (9.27 Å), b = 2.93 Å (2.86 Å), c = 4.60 Å (4.52 Å) (The values in the parenthesis are from the experiment³⁹). Layered NaMnO₂ is monoclinic (space group: C2/m) and optimized primitive lattice parameters a = 5.74 Å (5.67 Å), b = 2.92 Å (2.86 Å), c = 5.81 Å (5.81 Å) and β = 113.55°(113.2°) (The values in the parenthesis are from the experiment⁴⁰). 2a × 1b × 3c R-MnO₂ (Mn₂₄O₄₈) (Fig. 1a) and 4a × 1b × 3c layered NaMnO₂ supercells (Na₂₄Mn₂₄O₄₈) (Fig. 1b) are selected to explore the structural transformation of R-MnO₂ from the tunnel to the layered. To shorten the calculation time, the layered Na_12−xMn₁₂O₂₄ (0 ≤ x ≤ 12) system containing 48 atomic positions with a 1a × 3b × 2c supercell (Fig. 1c) is considered to study the electronic structures during charging.

Fig. 1 The crystal structure of the 2a × 1b × 3c R-MnO₂ (a), monoclinic layered NaMnO₂ structures with a 4a × 1b × 3c supercell (b) and a 1a × 3b × 2c supercell (c), the SEM image of R-MnO₂ (d), and XRD powder patterns (e) for layered NaMnO₂ and Na_0.33MnO₂ electrode after initial charge and discharge (solid lines) as well as calculated NaMnO₂ and Na_0.33MnO₂ (dots).

All the DFT calculations were performed using the Vienna Ab Initio Simulation package (VASP) within a projector augmented-wave (PAW) pseudo potential approach. A spin-polarized generalized gradient approximation (GGA) with the Perdew–Burke–Ernzerhof (PBE)⁴¹ exchange–correlation functional was employed. The electronic structures, Na 3s¹3p⁰, Mn 3d⁶4s¹ and O 2s²2p⁴, were selected. Considering that the element Mn has a higher angular quantum number of 3d electrons, we applied the GGA+U method described using the parameters U and J. The Hubbard U and J values of Mn are 6.0 eV and 1.0 eV, respectively, similar to the simplified U–J value of our previous⁴² and other works.⁴³ The energy cutoff for the plane waves was set to 520 eV. Monkhorst–Pack scheme 1 × 6 × 2 k-points meshes for 2a × 1b × 3c R-MnO₂ (Fig. 1a) and 4a × 1b × 3c layered NaMnO₂ supercells, and Monkhorst–Pack scheme 4 × 2 × 2 k-points meshes for 1a × 3b × 2c layered Na_12−xMn₁₂O₂₄ (0 ≤ x ≤ 12) were used for ion relaxations, lattice parameter optimization, and electronic structures. Atomic structures were fully optimized until forces acting on every atom became at least smaller than 0.02 eV Å⁻¹, and the total energy was converged to 10⁻⁴ eV.

The SEM image of R-MnO₂ is given in Fig. 1d. Secondary particles of R-MnO₂ show flower-like spherical morphology with a diameter of 1.7 μm, which is composed of rod-like crystallites with a width of about 50 nm. Simulated XRD from layered NaMnO₂ and Na_0.33MnO₂ fits with that of the initial discharging (NaMnO₂) and charging (Na_0.33MnO₂) material, respectively (Fig. 1e). These primary particles of R-MnO₂ less than 100 nm lead to the appearance of a large peak shape distribution. There appears a characteristic peak at 17° and a broad peak at 35–40°.

The phase structural transformation of MnO₂ from the tunnel to the layered under the different sodium contents is studied thermodynamically. The energy difference between both structures is expressed by ΔE:

ΔE = (E₁ − E₂)/2x (1)

E₁ and E₂ are the energies of Na_2xMn₂₄O₄₈ (0 ≤ x ≤ 12) systems with R-MnO₂ and layered NaMnO₂ structures, respectively. 2x is the number of Na atoms. When the absolute value of ΔE is obviously larger than 0, the tunnel structure or the layered structure is stable. However, when the absolute value of ΔE is close to 0, two structures might coexist, meaning that there are two phases during the first cycle discharge of R-MnO₂.

In order to understand the structural stability of layered Na_12−xMn₁₂O₂₄ (0 ≤ x ≤ 12) during the charge, the enthalpy change ΔH was calculated, and the corresponding formula was followed,

here, E(Na_12−xMn₁₂O₂₄), E(Na₁₂Mn₁₂O₂₄), and E(Mn₁₂O₂₄) are the total energies of layered Na_12−xMn₁₂O₂₄, layered Na₁₂Mn₁₂O₂₄, and layered Mn₁₂O₂₄ systems, respectively. x is the Na⁺ ion de-intercalation content. Structural distortion of layered NaMnO₂ derived from the continuous charge was described using distortion energies E_d:here, E(Mn₁₂O₂₄ − Na_12−xM₁₂O₂₄) and E(Mn₁₂O₂₄ − Na₁₂M₁₂O₂₄) are the total energies of Mn₁₂O₂₄ removing all Na⁺ ions from the layered Na_12−xMn₁₂O₂₄ and layered Na₁₂Mn₁₂O₂₄ systems, respectively.

3. Results and discussion

3.1. Geometrical structures and stabilities

Our experiment³⁷ found that the orthorhombic tunnel R-MnO₂ was transformed to the layered monoclinic NaMnO₂ during the first cycle discharge, then, the layered NaMnO₂ structure is retained, and the Na_0.33MnO₂ component is formed during the next cycle charge. In all, the change between the layered NaMnO₂ structure and the layered Na_0.33MnO₂ structure during the later cycle charge/discharge occurs. Therefore, in this work, the structural changes of R-MnO₂ during the discharge and layered NaMnO₂ during the charge and the structural transformation of MnO₂ from the tunnel to the layered are explored thermodynamically.

3.1.1. Structural change of R-MnO₂ during Na⁺ ion insertion. Various possible sites of Na⁺ ion insertion to orthorhombic R-MnO₂ are considered, and the system with the 4c (Fig. 2a) site has the lowest energy. When the number of Na⁺ ions inserted to Mn₂₄O₄₈ is more than 3, effects of mutual repulsion between Na⁺ ions open the 1 × 2 tunnels (Fig. 2c–e) until all 1 × 2 tunnels are opened as the NaMnO₂ component is formed (Fig. 2f). Li et al.³⁶ found that a common shearing and buckling movement of the MnO₂ layer led to the breaking of the Mn–O framework when the layered MnO₂ structure was transformed to MnO₂ ones with 1 × 1, 1 × 2, and 2 × 2 tunnels. In the same way, the breaking of the Mn–O framework will also occur when the MnO₂ structures with 1 × 1, 1 × 2, and 2 × 2 tunnels are transformed to the layered MnO₂ during Na⁺ ion insertion. The behavior similarly happens at the structural change of R-MnO₂ during Na⁺ ion insertion, which is consistent with Li's result.³⁶

Fig. 2 The tunnel Na_2xMn₂₄O₄₈ structures at 2x = 1 (a), 3 (b), 4 (c), 8 (d), 20 (e), and 24 (f).

3.1.2. Structures and stabilities of layered Na_12−xMn₁₂O₂₄ during Na⁺ ion de-intercalation. Stable geometrical structures of layered Na_12−xMn₁₂O₂₄ after the Na⁺ ion de-intercalation are obtained (Fig. 3). The de-intercalation rule of Na⁺ ions is to ensure that the remaining Na⁺ ions are uniformly dispersed. The lattice parameters are shown in Table 1.

Fig. 3 Structures of Na_12−xMn₁₂O₂₄ at x = 0 (a), 4 (b), and 8 (c) from the view point of the Na layer. ( Na O Mn).

Table 1 Calculated lattice parameters of the unit cell, the O–O layer spacing (d), volume contraction (ΔV), distortion energies (E_d), and the enthalpy change (ΔH) of Na_12−xMn₁₂O₂₄

x in Na12−xMn12O24	a (Å)	b (Å)	c (Å)	B (°)	d (Å)	ΔV (%)	E d (eV)	ΔH (eV)
0	5.74	2.92	5.81	113.55	3.20	—	—	0.07
1	5.64	2.93	5.80	112.45	3.68	0.90	0.0	−0.23
2	5.57	2.92	5.80	111.73	3.70	1.89	−0.62	−0.80
3	5.48	2.93	5.79	110.67	3.69	2.81	−1.17	−0.95
4	5.47	2.91	5.81	110.37	3.73	3.04	−0.95	−1.87
5	5.38	2.92	5.83	109.30	3.79	3.31	−1.05	−1.86
6	5.32	2.91	5.87	108.68	3.91	3.46	−1.01	−2.20
7	5.24	2.92	5.89	107.68	3.98	3.69	−1.04	−2.08
8	5.24	2.91	5.95	107.82	3.74	3.44	−0.91	−2.47
10	5.15	2.93	5.92	107.52	3.90	4.80	−0.88	−1.14
12	5.09	2.94	5.14	105.80	3.17	17.06	−0.80	0.0

---

From Table 1, it can be seen that β angle decreases as x increases. Lattice constant a takes on an obvious decreasing trend. Lattice constant c slightly increases. Lattice constant b is almost invariable. Due to the mutual Coulomb repulsion of oxygen ions, O–O layer spacing d between two different MnO₂ layers in Na_12−xMn₁₂O₂₄ increases in general as x increases. The unit cell volume contraction ΔV of Na_12−xMn₁₂O₂₄ compared to Na₁₂Mn₁₂O₂₄ increases with x. When all Na ions are removed, ΔV of 17.06% would be adverse, and the lattice c and the O–O layer spacing d for the system at x = 12 becomes small. Except for it, the largest value of the volume swell is only 4.80%. The changing features of volume contraction ΔV and the O–O layer spacing d in Na_0.33MnO₂ (Na_12−xMn₁₂O₂₄ at x = 8) are fluctuating, and it is interesting that ΔH of Na_0.33MnO₂ is the smallest, see Table 1.

To examine the thermodynamic stability of Na_12−xMn₁₂O₂₄ during the Na⁺ ion de-intercalation, the enthalpy change ΔH is calculated (Fig. 4a and Table 1). From the obtained ΔH, ΔH of Na_12−xMn₁₂O₂₄ at x = 8 is the lowest, indicating that the sodium-deficient phase could be experimentally synthesized. The distortion energies E_d derived from the Na⁺ ion de-intercalation are negative because of the shrinkage of Mn–O bonds and trend to a constant before x ≥ 3 (Fig. 4a and Table 1). In the layered NaMnO₂, Mn³⁺ with the d⁴ electron structure leads to Jahn–Teller distortion of MnO₆ octahedron, so there are two types of Mn–O bond length: short bonds A and long bonds B. As the de-intercalation amount x in Na_12−xMn₁₂O₂₄ increases, there are more and more Mn⁴⁺ until all Mn³⁺ are oxidized to Mn⁴⁺ at x = 12, so long bonds B shorten and their distribution range gradually converges to that of a short bond with 1.947 Å (Fig. 4b).

Fig. 4 The enthalpy change ΔH and distortion energies E_d (a) and Mn–O bond length (b) as x in Na_12−xMn₁₂O₂₄.

3.1.3. Structural transformation of R-MnO₂ to layered NaMnO₂. The phase structural transformation of MnO₂ from the tunnel to the layered under the different sodium contents is studied thermodynamically. The energy difference ΔE is plotted in Fig. 5a. The absolute value of ΔE is obviously larger than zero at 2x ≤ 1 (Na_0.04MnO₂) and 2x ≥ 10 (Na_0.42MnO₂) in Na_2xMn₂₄O₄₈ and the former ΔE is negative, and the latter ΔE is positive; however, the absolute value of ΔE is close to zero at 1 < 2x < 10. This means that Na_≤0.04MnO₂ is inclined to be the tunnel structure, Na_≥0.42−1MnO₂ is the layered structure, and Na_0.04−0.42MnO₂ is their mixed phase structure. The calculated result can be confirmed by the XRD experiment. The XRD powder pattern of discharged R-MnO₂ electrode is shown in Fig. 5b. It can be seen that the characteristic peaks of the R-MnO₂ phase at 2θ = 22.4° are gradually decreased as the sodium content increases. In Na_0.03MnO₂, there is a characteristic peak at 2θ = 22.4° and no characteristic peak at 2θ = 16.6°, and thus, Na_0.03MnO₂ is the R-MnO₂ phase with the tunnel structure. In Na_0.2MnO₂, the characteristic peak at 2θ = 16.6° appears, namely the layered monoclinic NaMnO₂ phase is formed. In Na_0.5MnO₂, the characteristic peaks of the R-MnO₂ phase at 2θ = 22.4° disappear, while the intensity of the characteristic peak with layered NaMnO₂ structure phase at 2θ = 16.6° increases, showing that only the layered NaMnO₂ phase exists in the electrode. The Na_0.2−0.5MnO₂ would be a mixed phase with the tunnel and the layered structures.

Fig. 5 Energy difference ΔE of Na_2xMn₂₄O₄₈ between R-MnO₂ and layered NaMnO₂ structures (a) and XRD powder patterns of Na_0.03MnO₂, Na_0.2MnO₂, Na_0.33MnO₂, Na_0.5MnO₂, Na_0.6MnO₂, and R-MnO₂ (expressed as initial) (b). The insert picture in (a) is O3-type phase structure of Na_0.33MnO₂, and share edges MnO₆ are denoted as “E”. Two XRD powder patterns at the bottom of (b) are those of R-MnO₂ and layered monoclinic NaMnO₂, respectively.

The layered Na_0.33MnO₂ has an O3-type monoclinic structure. It can be transformed from O′3-type phase structure by gliding of the MnO₂ slab without the breakage of Mn–O bonds. The characteristic of NaO₆ octahedron, sharing only edges in the O′3-type phase structure, also exists in O3-type, as shown in Fig. 5a.

3.2. Redox process of layered Na_12−xMn₁₂O₂₄

The density of state (DOS) and partial DOS (PDOS) directly reflecting the electronic structural change are calculated to understand the redox process of Mn in the layered Na_12−xMn₁₂O₂₄ system (Fig. 6). NaMnO₂ has a band gap of about 0.6 eV. It varies from being a semiconductor to a conductor with the sodium de-intercalation, and finally, the band gap is opened as all Na⁺ ions are removed (Fig. 6a). The Jahn–Teller distortion of Mn³⁺ leads to a split of two e_g orbits, among which d_z2 orbital is occupied at −1.0 to – 0 eV and d_x²−y² orbital is empty, as shown in Fig. 6b. In PDOS of Na_12−xMn₁₂O₂₄, all up-spin states of t_2g and one up-spin state of are occupied, and another up-spin state of and both down-spin states of t_2g and are unoccupied in Na_12-xMn₁₂O₂₄ at x = 0, inferring Mn ions are in the +3 state. With the sodium de-intercalation, more and more up-spin d_z² states are unoccupied, indicating the Mn⁴⁺ ion formation (Fig. 6b). This can be confirmed by the XPS spectra of the R-MnO₂ electrode after the 2nd charge.³⁷ In the XPS spectra, Mn³⁺ and Mn⁴⁺ characteristic peaks are located at ∼640.5 eV and ∼641.8 eV, respectively, as in our previous work.³⁷ Oxidization of Mn³⁺ in Na_12−xMn₁₂O₂₄ is also supported by Bader charges, as shown in Table 2. The Bader charge of Mn³⁺ is 5.358 e − 5.411 e, and the Bader charge of Mn⁴⁺ is 5.272 e − 5.351 e.

Fig. 6 Evolution of TDOS (a) and PDOS of Mn³⁺/Mn⁴⁺ (b) for Na_12−xMn₁₂O₂₄.

Table 2 Bader charges of Mn⁴⁺ and Mn³⁺ in layered Na_12−xMn₁₂O₂₄. The number in parenthesis is that of Mn⁴⁺ or Mn³⁺

x	0	1	2	3	4	5	6	7	8	10	12
Mn^4+	—	5.351 (1)	5.310 (2)	5.315 (3)	5.288 (4)	5.302 (5)	5.289 (6)	5.295 (7)	5.287 (8)	5.277 (10)	5.272 (12)
Mn^3+	5.411 (12)	5.405 (11)	5.399 (10)	5.389 (9)	5.389 (8)	5.379 (7)	5.375 (6)	5.369 (5)	5.361 (4)	5.358 (2)	—

---

In order to visually understand the Mn⁴⁺ ion formation, the local charge density of Mn ions in the Na_0.33MnO₂ system located at −1.0 to –0 eV is presented, see Fig. 7. From Fig. 7a, it can be seen that electron density only focuses on 4 Mn ions among 12 Mn ions (0.33 ratio), namely 0.67 Mn³⁺ ions in Na_0.33MnO₂ are oxidized to Mn⁴⁺. We can also see that Mn³⁺ 3d_z2 and O 2p_z form a σ* anti-bond (The corresponding Mn³⁺–O bond length is 2.214 Å, see Fig. 8d in the following part), consistent with the charge density difference result (Fig. 7b). In Fig. 7b, the Mn⁴⁺–O bond is obviously ionic, while the Mn³⁺–O bond has anti-bonding characteristics due to charge repulsion between the positive charge density from Mn³⁺ and O²⁻.

Fig. 7 Charge density at –1.0 to –0 eV (a) and the right of (a) is the figure of the interaction between Mn³⁺ and O²⁻ ions. Charge density difference of Na_0.33MnO₂ (b) and the right of (b) gives a slice of charge density difference.

Fig. 8 The 1st–2nd voltage capacity profile of R-MnO₂ in the EC/PC electrolyte (a). The relative Fermi energy level of layered Na_12−xMnO₂ and R-MnO₂ (b), where the Fermi level of NaMnO₂ is set to 0. Energy band structures and HOMO of NaMnO₂ (c), Na_0.33MnO₂ (d), and R-MnO₂ (e). Here, Mn–O bond lengths of one MnO₆ octahedron are given, and the unit is Å.

3.3. Potential plateau during the discharge

From our experiment,³⁷ it is known that R-MnO₂ forms the layered NaMnO₂ after the initial discharge, NaMnO₂ retains its layered structure after the initial charge, and the next discharge/charge is the transformation between the layered Na_0.33MnO₂ and the layered NaMnO₂. The initial activation process and the structural change during the charging/discharging are similar to other Mn-based materials such as Li[Li_0.2Ni_0.16Mn_0.56Co_0.08]O₂.³⁸ Potential plateaus of R-MnO₂ turning into the layered NaMnO₂ and the layered Na_0.33MnO₂ turning into the layered NaMnO₂ are different (Fig. 8a). In addition, the complete charge–discharge curve and cycling performance have been taken in our previous work using BMIMTFSI and EC/PC electrolytes.³⁷ In order to explain the phenomenon, Fermi energy levels of various layered Na_12−xMn₁₂O₂₄ compositions and R-MnO₂ are calculated (Fig. 8b). The aim is to use two differences between Fermi energy levels of Na_12−xMn₁₂O₂₄ and NaMnO₂, and R-MnO₂ and NaMnO₂ to reflect two different potential plateaus of transformations from Na_12−xMn₁₂O₂₄ to NaMnO₂ and R-MnO₂ to NaMnO₂ during the discharge. Fermi energy of one bulk system can be obtained by calculating the Fermi energy of its most stable surface or its first ionization potential (IP). Our work selects the second method. IP is the difference between total energies of the neutral substance and the system with one electron removed to infinity, expressed as

IP = E(+) − E(0) (4)

E(+) and E(0) are the energies per Na_(12−x)/12MnO₂ formula with and without one positive charge, respectively. In Fig. 8b, the Fermi energy level of NaMnO₂ is the lowest, it lifts with the increase in x of Na_12−xMn₁₂O₂₄, and the Fermi energy levels of R-MnO₂ and Na₁₁Mn₁₂O₂₄ are close. The Fermi energy difference of R-MnO₂ to NaMnO₂ is lower than that of Na_(12−x)/12MnO₂ to NaMnO₂. So, the phenomenon that the potential plateau of R-MnO₂ turning into the layered NaMnO₂ is lower than that of the layered Na_0.33MnO₂ transforming to the layered NaMnO₂ is seen in our experiment result.

The HOMO and the geometrical structures of the layered Na_0.33MnO₂, layered NaMnO₂, and R-MnO₂ are employed to understand their Fermi energy level variation. To obtain HOMO, energy band structures of these systems are first calculated to get the highest occupied energy level at a certain k-point, and subsequently, the electronic structure at the energy level, namely HOMO, is computed (Fig. 8c–e). The MnO₆ octahedron of NaMnO₂ has two long Mn–O bonds (2.428 Å) and four short Mn–O bonds (1.969 Å) due to the Jahn–Teller distortion of Mn³⁺3d⁴. In HOMO of NaMnO₂, the interaction between O 2p and Mn 3d in four short Mn–O bonds is non-bonding. There is a similar case between O 2p and Mn 3d in three shorter Mn–O bonds (1.843 Å, 1.843 Å, and 1.871 Å) for R-MnO₂. In NaMnO₂, the Jahn–Teller distortion of Mn³⁺ 3d⁴ would benefit the structural stability. So, the Fermi energy level of NaMnO₂ is lower than that of R-MnO₂. The Na_0.33MnO₂ system has Mn³⁺ and Mn⁴⁺. The interaction between O 2p and Mn 3d in the long Mn³⁺–O bonds (2.214 Å) is anti-bonding because the direction of O 2p orbital and Mn³⁺ 3d orbital is in line. This would lead to a higher Fermi energy level of Na_0.33MnO₂ than NaMnO₂ and R-MnO₂ systems. The interaction between O 2p and Mn 3d in HOMO of other Na_(12−x)/12MnO₂ systems have the same anti-bonding characteristic to Na_0.33MnO₂, resulting in a whole rule that the Fermi energy level difference between R-MnO₂ and NaMnO₂ is lower than that between Na_(12−x)/12MnO₂ and NaMnO₂.

4. Conclusion

In this work, nanosized flower-like R-MnO₂ with a 1 × 2 tunnel prepared experimentally is discharged to obtain a monoclinic layered NaMnO₂ cathode material. The next cycle charge/discharge transformation happens between the two layered structures of NaMnO₂ and Na_0.33MnO₂. During the charging process, the cathode material delivers 206 MA h g⁻¹ through the redox couple of Mn³⁺/Mn⁴⁺. R-MnO₂ retains its structure when the molar content of inserted Na does not exceed the value of 0.04. It is transformed into the layered NaMnO₂ structure when the insertion Na molar content is ≥0.42. The middle composition is the mixed phase of the tunnel and the layered structures. The different potential plateaus of R-MnO₂ turning into the layered NaMnO₂ and the layered Na_0.33MnO₂ transforming to the layered NaMnO₂ are from the distinctions of two Fermi level differences between R-MnO₂/NaMnO₂ and Na_0.33MnO₂/NaMnO₂. The work not only provides new insight into a deeper understanding of the redox chemistry in sodium ion cathode materials, but also facilitates the development of higher capacity electrons for sodium ion batteries.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

The authors appreciate financial support from China-US Electric Vehicle Project (S2016G9004), National Key Project on Fundamental Research Program (2015CB251104), and Inner Mongolia Natural Science Foundation (No. 2021MS02003).

References

1. N. OrtizVitoriano, E. Drewett, E. Gonzaloa and T. Rojo, High performance manganese-based layered oxide cathodes: overcoming the challenges of sodium ion batteries., Energy Environ. Sci., 2017, 10, 1051–1074.
2. S. Kim, D. H. Seo, X. Ma, G. Ceder and K. Kang, Electrode Materials for Rechargeable Sodium–Ion Batteries: Potential Alternatives to Current Lithium–Ion Batteries, Adv. Energy. Mater., 2012, 2, 710–721.
3. R. S. Carmichael, Practical handbook of physical properties of rocks and minerals, CRC Press, 1989.
4. D. P. K. Nayak, L. Yang, W. Brehm and P. Adelhelm, From Lithium–Ion to Sodium–Ion Batteries: Advantages, Challenges, and Surprises, Angew. Chem., Int. Ed., 2018, 57, 102.
5. J. Zhang, Y. Fang, L. Xiao, J. Qian, Y. Cao, X. Ai and H. Yang, From Lithium–Ion to Sodium–Ion Batteries: Advantages, Challenges, and Surprises, ACS Appl. Mater. Interfaces, 2017, 9, 7177.
6. X. Wu, G. Zhong and Y. Yang, Sol–gel synthesis of Na₄Fe₃(PO₄)₂(P₂O₇)/C nanocomposite for sodium ion batteries and new insights into microstructural evolution during sodium extraction, J. Power Sources, 2016, 327, 666–674.
7. Z. Jian, L. Zhao, H. Pan, Y. S. Hu, H. Li, W. Chen and L. Chen, Carbon coated Na₃V₂(PO₄)₃ as novel electrode material for sodium ion batteries, Electrochem. Commun., 2012, 14, 86–89.
8. Z. Chen, T. Yuan., X. Pu, H. Yang, X. Ai, Y. Xia and Y. Cao, Symmetric Sodium–Ion Capacitor Based on Na_0.44MnO₂ Nanorods for Low-Cost and High-Performance Energy Storage, ACS Appl. Mater. Interfaces, 2018, 10, 11689–11698.
9. M. Xie, R. Luo, J. Lu, R. Chen, F. Wu, X. Wang, C. Zhan, H. Wu, H. M. Albishri and A. S. Al-Bogami, Synthesis-Microstructure-Performance Relationship of Layered Transition Metal Oxides as Cathode for Rechargeable Sodium Batteries Prepared by High-Temperature Calcination, ACS Appl. Mater. Interfaces, 2014, 6, 17176–17183.
10. K. Kubota, N. Yabuuchi, H. Yoshida, M. Dahbi and S. Komaba, Layered oxides as positive electrode materials for Na-ion batteries, MRS Bull., 2014, 39, 416–422.
11. M. H. Han, E. Gonzalo, G. Singh and T. Rojo., A comprehensive review of sodium layered oxides: powerful cathodes for Na-ion batteries, Energy Environ. Sci., 2015, 8, 81–102.
12. N. Yabuuchi, M. Kajiyama, J. Iwatate, H. Nishikawa, S. Hitomi, R. Okuyama, R. Usui, Y. Yamada and S. Komaba, P2-type Na_x[Fe_1/2Mn_1/2]O₂ made from earth-abundant elements for rechargeable Na batteries, Nat. Mater., 2012, 11, 512–517.
13. J. P. Parant, R. Olazcuaga and M. Devalette, Sur Quelques Nouvelles Phases de Formule Na_xMnO₂(x = 1), J. Solid State Chem., 1971, 3, 1–11.
14. J. Billaud, G. Singh, A. R. Armstrong, E. Gonzalo, V. Roddatis, M. Armand, T. Rojo and P. G. Bruce, Na_0.67Mn_1−xMg_xO₂ (0 ≤ x ≤ 0.2): a high capacity cathode for sodium-ion batteries, Energy Environ. Sci., 2014, 7, 1387–1391.
15. N. Sharma, N. Tapia-Ruiz, G. Singh, A. R. Armstrong, J. C. Pramudita, H. E. A. Brand, J. Billaud, P. G. Bruce and T. Rojo, Rate Dependent Performance Related to Crystal Structure Evolution of Na_0.67Mn_0.8Mg_0.2O₂ in a Sodium–Ion Battery, Chem. Mater., 2015, 27, 6976–6986.
16. J. S. Thorne., R. A. Dunlap and M. N. Obrovac, A Combinatorial study of the Sn–Si–C system for Li-ion battery applications, J. Electrochem. Soc., 2012, 160, A361–A367.
17. X. Wang, M. Tamaru, M. Okubo and A. Yamada, Electrode Properties of P2-Na_2/3Mn_yCo_1−yO₂ as Cathode Materials for Sodium–Ion Batteries, J. Phys. Chem. C, 2013, 117, 15545–15551.
18. Z. Lu and J. R. Dahn, Situ X-Ray Diffraction Study of P2 Na_2/3[Ni_1/3Mn_2/3]O₂, J. Electrochem. Soc., 2001, 148, A1225–A1229.
19. I. Hasa, D. Buchholz, S. Passerini, B. Scrosati and J. Hassoun, High Performance Na0.5[Ni0.23Fe0.13Mn0.63]O2 Cathode for Sodium-Ion Batteries, Adv. Energy Mater., 2014, 4, 1400083.
20. M. H. Han, E. Gonzalo, N. Sharma, J. M. Lopez Del Amo, M. Armand, M. Avdeev, J. J. Saiz, J. J. S. Garitaonandia and T. Rojo, High-Performance P2-Phase Na_2/3Mn_0.8Fe_0.1Ti_0.1O₂ Cathode Material for Ambient-Temperature Sodium-Ion Batterie, Chem. Mater., 2016, 28, 106–116.
21. D. Yuan, W. He, F. Pei, F. Wu, Y. Wu, J. Qian, Y. Cao, X. Ai and H. Yang, Synthesis and electrochemical behaviors of layered Na_0.67[Mn_0.65Co_0.2Ni_0.15]O₂ microflakes as a stable cathode material for sodium-ion batteries, J. Mater. Chem. A, 2013, 1, 3895–3899.
22. L. Zheng, J. Li and M. N. Obrovac, Crystal Structures and Electrochemical Performance of Air-Stable Na_2/3Ni_1/3−x,Cu_xMn_2/3O₂ in Sodium Cells, Chem. Mater., 2017, 29, 1623–1631.
23. H. Yoshida, N. Yabuuchi, K. Kubota, I. Ikeuchi, A. Garsuch, M. Schulz-Dobrick and S. Komaba, P2-type Na_2/3Ni_1/3Mn_2/3−xTi_xO₂ as a new positive electrode for higher energy Na-ion batteries, Chem. Commun., 2014, 50, 3677–3680.
24. X. H. Zhu, F. Q. Meng, Q. H. Zhang, L. Xue, H. Zhu, S. Lan, Q. Liu, J. Zhao, Y. H. Zhuang, Q. B. Guo, B. Liu, L. Gu, X. Lu, Y. Ren and H. Xia, LiMnO₂ Cathode Stabilized by Interfacial Orbital Ordering for Sustainable Lithium-Ion Batteries, Nat. Sustain., 2021, 4, 392–401.
25. X. H. Zhu, Q. Y. Xia, X. Y. Liu, Q. H. Zhang, J. Xu, B. W. Lin, S. Li, Y. H. Zhuang, C. Qiu, L. Xue, L. Gu and H. Xia, Retarded Layered-to-Spinel Phase Transition in Structure Reinforced Birnessite with High Li Content, Sci. Bull., 2021, 66, 219–224.
26. N. Wang, H. Pang, H. Peng, G. Li and X. Chen, Hydrothermal synthesis and electrochemical properties of MnO₂ nanostructures, Cryst. Res. Technol., 2009, 44, 1230–1234.
27. C. M. Julien and A. Mauger, Nanostructured MnO₂ as Electrode Materials for Energy Storage, Nanomaterials, 2017, 7, 396.
28. Y. Zhou, T. Chen, J. Zhang, Y. Liu and P. Ren, Amorphous MnO₂ as Cathode Material for Sodium-ion Batteries, Chin. J. Chem., 2017, 35, 1294–1298.
29. D. Lv, X. Huang, H. Yue and Y. Yang, Sodium-Ion-Assisted Hydrothermal Synthesis of gamma-MnO₂ and Its Electrochemical Performance, J. Electrochem. Soc., 2009, 156, A911–A916.
30. A. T. David. and M. S. Islam, Electrochemistry of Hollandite alpha-MnO₂: Li–Ion and Na–Ion Insertion and Li₂O Incorporation, Chem. Mater., 2013, 25, 2515–2526.
31. Y. An, J. Feng, L. Ci and S. Xiong, MnO₂ nanotubes with a water soluble binder as high performance sodium storage materials, RSC Adv., 2016, 6, 103579.
32. B. Kishore, G. Venkatesh and N. Munichandraiah, A Na/MnO₂ Primary Cell Employing Poorly Crystalline MnO₂, J. Electrochem. Soc., 2015, 162, A839–A844.
33. D. Su, H. J. Ahn and G. Wang, Beta-MnO₂ nanorods with exposed tunnel structures as high-performance cathode materials for sodium-ion batteries, NPG. Asia Mater., 2013, 5, e70.
34. D. Su, H. J. Ahn and G. Wang, Hydrothermal synthesis of alpha-MnO₂ and beta-MnO₂ nanorods as high capacity cathode materials for sodium ion batteries, J. Mater. Chem. A, 2013, 1, 4845.
35. Y. Liu, Y. Qiao, W. Zhang, H. Wang, K. Chen, H. Zhu, Z. Li and Y. Huang, Nanostructured alkali cation incorporated delta-MnO₂ cathode materials for aqueous sodium-ion batteries, J. Mater. Chem. A, 2015, 3, 7780–7785.
36. Y. F. Li, S. C. Zhu and Z. P. Liu, Reaction Network of Layer-to-Tunnel Transition of MnO₂, J. Am. Chem. Soc., 2016, 138, 5371–5379.
37. H. N. Si, L. Li, W. J. Hao, L. Seidl, X. L. Cheng, H. Y. Xu, G. X. Jia, O. Schneider, S. L. An and X. P. Qiu, Structural Transformation and Cycling Improvement of Nanosized Flower-like γ-MnO₂ in a Sodium Battery, ACS. Appl. Energy. Mater., 2019, 2, 5050–5056.
38. D. Buchholz, J. Li, S. Passerini, G. Aquilanti, D. D. Wang and M. Giorgetti, X-Ray Absorption Spectroscopy Investigation of Lithium-Rich, Cobalt-Poor Layered-Oxide Cathode Material with High Capacity, ChemElecctroChem., 2015, 2, 85–97.
39. J. E. Post and P. J. Heaney, Neutron and synchrotron X-ray diffraction study of the structures and dehydration behaviors of ramsdellite and “groutellite”, Am. Mineral., 2004, 89, 969–975.
40. O. I. Velikokhatnyi, D. Choi and P. N. Kumta, Energetics of the lithium-magnesium imide–magnesium amide and lithium hydride reaction for hydrogen storage: An ab initio study, Mater. Sci. Eng., B, 2006, 128, 115–124.
41. J. P. Perdew, K. Burke and M. Ernzerhof, Generalized gradient approximation made simple, Phys. Rev. Lett., 1996, 77, 3865–3868.
42. H. Z. Wei, X. L. Cheng, H. W. Fan, Q. Shan, S. L. An, X. P. Qiu and G. X. Jia, A Cobalt-Free Li(Li_0.17Ni_0.17Fe_0.17Mn_0.49)O₂ Cathode with More Oxygen-Involving Charge Compensation for Lithium-Ion Batteries, ChemSusChem, 2019, 12, 2471–2479.
43. Z. Z. Zhang, Y. F. Zhang, Y. Li, J. Lin, D. G. Truhlar and S. P. Huang, MnSb₂S₄ Monolayer as an Anode Material for Metal-Ion Batteries, Chem. Mater., 2018, 30, 3028–3214.

---

Footnote

† These authors contribute equally to this work.

---