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

Kai Qiu† ^a, Chao Zhang† ^abc, Mingxia Yan ^a, Shouwang Zhao ^a, Hongwei Fan ^a, Shengli An *^a, Xinping Qiu *^b and Guixiao Jia *^a
aSchool of Materials and Metallurgy, Inner Mongolia University of Science and Technology, Baotou, Inner Mongolia 014010, China. E-mail: shengli_an@126.com; guixiao.jia@163.com
^bKey Laboratory of Organic Optoelectronics and Molecular Engineering, Department of Chemistry, Tsinghua University, 100084, China. E-mail: qiuxp@mail.tsinghua.edu.cn
^cSchool of Metallurgical and Ecological Engineering, University of Science and Technology Beijing, 100083, China

First published on 23rd November 2021

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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.

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,

	(2)

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:

	(3)

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.

3.2. 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.

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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⁻.

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).

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Footnote

† These authors contribute equally to this work.

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