Facile synthesis and performance of Na-doped porous lithium-rich cathodes for lithium ion batteries

Abstract

Na-doped porous lithium-rich (Li-rich) cathode microspheres (∼1 μm) were firstly prepared via the solvothermal method and subsequently a high-temperature calcination process. X-ray diffraction (XRD), scanning electron microscopy (SEM), transmission electron microscopy (TEM), X-ray photoelectron spectroscopy (XPS), and nitrogen adsorption–desorption isotherms were used to characterize the structure and morphology of the as-prepared cathode material. It has been found that the as-prepared material has an obvious internal porous structure with the existence of 5 at% Na ions. Besides, the obtained cathode material possesses excellent electrochemical characteristics. For example, it can deliver a high initial discharge capacity of 305.2 mA h g⁻¹ between 2.0 V and 4.6 V at a rate of 0.1C at room temperature and the retention of the capacity is still as high as 88.2% after 200 cycles. Furthermore, the electrochemical impedance spectroscopy (EIS) results also show that the introduction of Na ions can decrease the charge transfer resistance of the as-prepared cathode material. The excellent electrochemical performance of the as-prepared cathode material can probably be attributed to the improved stability of the bulk lattice and the expanded Li slab space, which facilitates lithium ion diffusion and effectively enhances the stability of the layered structure of the materials.

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1. Introduction

During the past decades, Li-rich cathodes have been intensely investigated as a promising cathode material for Li-ion batteries due to their high operational voltage of 4.6 V and high practical capacity of 200–260 mA h g⁻¹.^1–3 Among various candidates, Li-rich Mn-based oxide materials, usually denoted as xLi₂MnO₃·(1 − x)LiMO₂ or Li_1+x/(x+2)M_{(2−2x)/(x+2)}Mn_2x/(x+2)O₂ (M = Mn, Ni, Co, etc.), have received particular attention due to their high reversible capacity (∼300 mA h g⁻¹), which is nearly twice as high as those of the LiCoO₂ and LiFePO₄ cathodes presently used in commercial LIBs.⁴ Many recent investigations have also indicated that Li-rich Mn-based oxides could be highly promising cathode materials for advanced Li-ion batteries with potential use in all electric vehicles because of their extremely high energy density.⁵ However, despite these merits and application advantages, Li-rich Mn-based oxide materials have several drawbacks, including a low initial columbic efficiency, poor rate capability, and insufficient capacity retention, which hinder their practical applications.^3–5 Generally speaking, the low initial coulombic efficiency of these materials is usually considered to be caused by the elimination of the oxygen vacancies formed during the first charge due to the irreversible removal of Li₂O from the crystal lattice. The removal of Li₂O also causes a rearrangement of the transition metal ions in the crystal lattice, which leads to structural instability and thus causes capacity degradation during cycling; meanwhile, the formation of a solid electrolyte interface (SEI) film may slow down the kinetics of the lithium ion extraction/insertion in the surface region, and in addition, the transport behavior in the bulk phase is also considered to be important for the rate capability.^6,7

Considering the fact that the drawbacks of these materials can be adjusted efficiently through controlling their structures, several strategies have been successfully developed to fabricate materials with different morphologies and desired micro- or nanostructures. One effective approach is to modify the surface of the cathode active materials with conductive polymers as well as inorganic materials such as poly(3,4-ethylenedioxythiophene) (PEDOT), oxides (Al₂O₃, TiO₂), phosphates (AlPO₄, CoPO₄), and fluorides (AlF₃).^8–10 The coating layers can really alleviate the elimination of the oxygen vacancies and the decomposition of the electrolyte so as to decrease the initial irreversible capacity and effectively prevent attack of the surface of the electrode by HF in the electrolyte.¹¹ However, this beneficial effect may be contrasted with an increase of the battery impedance resulting from the insulating nature of coating media in segmental materials. In addition, the coating layer is usually so thin and heterogeneous that it can wear off the cathode surface during long-term cycling if it is subjected to continuous HF attack, leading to a slow fading of the capacity during extensive cycling.^12,13 Besides, it has been reported that shortening the Li ion migration length may be a more effective way to improve the rate performance. This also triggers investigations about nano-modification technology for cathode materials in order to improve the rate performance of Li-ion batteries. Compared with the secondary particle size, the primary particle size usually has a greater influence on the electrochemical and physical properties.^12–15 However, for Li-rich materials, if the particle size decreases to the nanoscale, side-reactions will be promoted due to the lower packing density and higher specific surface area of the nanoparticles. It is not necessarily true that a smaller particle size will certainly achieve a better electrochemical performance.^15,16 A hollow structure is another interesting method to improve the electrochemical performances of the cathode materials. Firstly, the cavity in the hollow structures can provide extra active sites for Li⁺ storage, which is beneficial to enhance the specific capacity. Secondly, this type of structure can provide a shorter path for Li⁺ ion diffusion and meanwhile has a larger electrode–electrolyte contract area for Li⁺ flux across the interface, resulting in an improvement of the rate capacity.^17–19 However, the preparation of a hollow structure for the micro-particles requires a more complicated design process and control model. Therefore, it is still a great challenge to form hollow structures for Li-rich metal oxides while controlling the interiors, despite its popularity in simple metal oxides.²⁰ Compared with a hollow structure, a porous structure is a much better approach for the improvement of the electrochemical performance due to its higher packing density and better structural stability based on the void space within each sphere. It is believed that the structural advantages of the porous microsphere may include the high packing density and the decrease of electrode polarization in the active layer.¹⁰ These advantages can provide the enhanced specific capacities that originate from the extra active sites used for the storage of Li⁺ within the holes. Furthermore, based on the conception of the microsphere or nanosphere, preparation and structural control of the porous structure are easier than for a hollow structure. In addition to the methods listed above, element doping is another effective method to tune the structure and performances of electrode materials. For example, a better capacity retention and higher discharge capacity for LiMnO₂ have been achieved after Zn, Fe, Co, Cr, or Al doping, and improved cycling stability, lithium diffusivity, and electronic conductivity have also been observed after Cr or Ru doping of Li₂MnO₃.^21–24

In consideration of the advantages and disadvantages of the above mentioned strategies, it is vitally important to choose a befitting preparation method to improve more effectively the electrochemical performance of materials. In our previously reported works,^25–27 a sphere Li-rich layered material with an average diameter of 1 μm was synthesized via the solvothermal method. Furthermore, although doping with Na ions has been reported as an effective method to improve the electrochemical performance of cathode materials,²⁸ its combination with a porous structure has rarely been reported, which can not only meet the requirements of controllability and stability of the interior structure by the introduction of Na ions into the porous structure, but can also enhance the electrochemical conductivity and cycling stability by the freer Li⁺ insertion/extraction processes. Besides, there is a large ionic radii difference between the Na ion (r_Na⁺ = 1.02 Å) and the transition metal ions (Ni²⁺, Mn⁴⁺, or Co³⁺), and a larger driving force can be generated to separate the alkali ions from the transition metal layer thus avoiding substantial alkali ion/transition metal ion disorder. At the same time, introducing Na ions into the Li layer can stabilize the bulk lattice and expand the Li slab space so as to facilitate lithium ion diffusion and effectively enhance the stability of the layered structure.^10,28 Comparing with a hollow structure and a concentration-gradient structure, a porous structure is considered to be a feasible and befitting design for this microsphere because this structure can keep the structural integrity via a moderate and controllable method, and the porous structure can provide extra active sites for Li⁺ storage in the meantime.²⁹ Thus, in order to further improve the electrochemical performance of the as-prepared Li-rich cathode material, e.g., the cycle stability, rate performance, and initial coulombic efficiency, this work is mainly focused on the design and preparation of a Na-doped Li-rich microsphere material with a porous structure distributed inside of the spheres.

2. Experimental

2.1. Synthesis of the Na-doped Li_1.4[Mn_0.6Ni_0.2Co_0.2]O₂ porous microspheres

All reagents were of analytical grade and were used without further purification. The porous cathode microspheres were firstly synthesized via the solvothermal method and a subsequent high-temperature solid state method. For example, the Mn_0.6Ni_0.2Co_0.2CO₃ microspheres were first synthesized via the solvothermal method of MnCl₂·4H₂O, CoCl₂·6H₂O and Ni(NO₃)₂·6H₂O (3 : 1 : 1 in molar ratio) as starting materials. These powers were dissolved into 50 mL ethanediol under magnetic stirring. Then, the prepared NH₄HCO₃ solution was added dropwise into the mixed solution. This resultant mixture was continually stirred for 30 min to obtain a homogeneous solution and then transferred into a Teflon lined stainless-steel autoclave (capacity of 100 mL). Finally, the autoclave was sealed and maintained at 200 °C for 20 h in an electron oven. The system was then cooled to ambient temperature naturally. The obtained precipitate was centrifugalized and washed with ethanol several times, followed by vacuum-drying at 60 °C. This prepared Mn_0.6Ni_0.2Co_0.2CO₃ precursor was then ready for further processing and characterization.

To attain conversion from the Mn_0.6Ni_0.2Co_0.2CO₃ microsphere precursor to the [Mn_0.6Ni_0.2Co_0.2]O₄ porous microsphere, thermal treatment was performed at 400 °C for 10 h with a heating ramp of 4 °C min⁻¹ until the products changed color from pink to black. To obtain the Na-doped Li_1.4[Mn_0.6Ni_0.2Co_0.2]O₂ porous microspheres, the prepared [Mn_0.6Ni_0.2Co_0.2]O₄ microspheres (2.2 g) were dispersed in absolute ethanol (50 mL) and mixed with an appropriate amount of NaCO₃ under ultrasound conditions for 30 min. After that, the mixtures were thoroughly mixed with an appropriate amount of lithium carbonate, and the mixture was preheated at 500 °C for 8 h, and then calcined at 900 °C for 12 h in air to form the final Na-doped Li_1.4[Mn_0.6Ni_0.2Co_0.2]O₂ porous cathode material, which was named Na-LMO. In order to make a comparison, a Li_1.4[Mn_0.6Ni_0.2Co_0.2]O₂ cathode microsphere was synthesized in a similar way to the above, expect the thermal treatment was performed at 400 °C and it was not mixed with Na₂CO₃. The prepared sample was named LMO to distinguish it from the previous Na-LMO sample.

2.2. Characterization of morphology and structure

The phase identification of the samples was performed with a diffractometer (D/Max-3C, Rigaku, Japan) using Cu Kα radiation (λ = 1.54178 Å) and a graphite monochromator at 36 kV, 20 mA. The scanning rate was 2° min⁻¹, and the scanning range of the diffraction angle (2θ) was 10° ≤ 2θ ≤ 80°. The morphology and atomic concentration of the samples were observed by scanning electron microscopy (SEM) and energy-dispersive X-ray spectroscopy (EDS) (JSM-6100LV, JEOL, Japan). The chemical compositions of the as-obtained materials were determined using atomic absorption spectroscopy (AAS, Vario 6 Analytik Jena AG, Jena) measurements. Transmission electron microscopy (TEM, FEITecnai G2, Holland) was used to investigate the crystal structure of the materials. The surface solid-state chemistry of the particles was characterized by X-ray photoelectron spectroscopy (XPS, KAlpha 1063, Thermo Fisher Scientific, England). The activated nitrogen-enriched novel carbons (a-NENCs) were examined by adsorption experiments of nitrogen and adsorption/desorption isotherms of nitrogen were measured at 77 K on a Quantachrome autosorb automated gas sorption system. The estimation of the Brunauer–Emmett–Teller (BET) specific surface area, pore volume, and pore size distribution (PSD) were carried out according to the BET equation and Barrett–Joyner–Halenda (BJH) method.

2.3. Electrochemical measurements

To prepare the positive electrodes, a mixture of the as-prepared cathode materials (80 wt%), polyvinylidene fluoride (10 wt%), graphite (5 wt%), and acetylene black (5 wt%), dissolved in a N-methylpyrrolidone (NMP) solution, was coated onto Al foil. Generally, the loading mass of the active substances was 4–5 mg cm⁻². The coated Al foil was dried at 110 °C for 24 h in a vacuum oven before processing. Then the positive electrodes were fabricated by cutting the coated Al foil into circular films with 10 mm diameters. CR2025 coin cells were assembled in a glove box using lithium foil as the anode, 1 mol L⁻¹ LiPF₆ in ethylene carbonate (EC)–dimethyl carbonate (DMC) (1 : 1, v/v) as the electrolyte, and a porous polypropylene based membrane (Celgard) as the separator, under argon atmosphere. To determine the electrochemical characteristics of the as-obtained cathode materials, the cells were charged and discharged at different current densities between 2.0 V and 4.6 V (vs. Li⁺/Li), using a Neware battery test system (Shenzhen, China).

3. Results and discussion

As shown in Fig. 1, the morphology and microstructure of the as-prepared microspheres were characterized by SEM. The typical low magnification SEM images in Fig. 1(a) and (b) show that the carbonate precursor and oxide precursor particles obtained from the solvothermal process and calcination process have a spherical morphology, the shape of the particle is regular, and the size is homogeneously distributed in the range of 1–2 μm. Meanwhile, in the images in Fig. 1(c) and (d), the Na-LMO particles appear to be homogeneous spheres with a size distribution of 1–2 μm, retaining the unified size with previous particles. It should be noted that the resultant samples in Fig. 1(c) and (d) show that the particle is composed of uniform primary grains, and a series of holes can be observed from Fig. 1(c) and (d). Obviously, it can be seen from Fig. 1 that the morphologies of the as-prepared product particles and precursor particles are very different, which is possibly because the introduction of Na ions into the Li slab can improve the stability of the layer structure so that the growth of the crystal tends to follow the direction of the layer.¹⁰ As a further quantitative analysis of the Na-LMO sample, the samples Na-LMO and LMO were examined via EDS (in Fig. 2) and atomic absorption spectroscopy (AAS). The results in Fig. 2 and Table 1 show that about 5 at% Na can be detected in the Na-LMO sample and the transition metal atom relative ratios have no evident difference between these two samples, which is in accordance with the theoretical ratios (Mn : Ni : Co = 6 : 2 : 2).

Fig. 1 SEM images of the (a) carbonate precursor and (b) oxide precursor, (c) and (d) Na-LMO sample.

Fig. 2 EDS spectra of the LMO and Na-LMO samples.

Table 1 The relative amounts of Li, Ni, Co, and Mn in Na-LMO and LMO

Sample	Li	Ni	Co	Mn	Rate of Li/Ni/Co/Mn
Na-LMO	1.505	0.205	0.200	0.592	Li1.5Ni0.2Co0.2Mn0.6
LMO	1.497	0.204	0.200	0.592	Li1.5Ni0.2Co0.2Mn0.6

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To elucidate the accurate morphologies and detailed crystal structures of the final prepared materials, the results of TEM, high-resolution TEM (HRTEM) and selected area electron diffraction (SAED) studies are shown in Fig. 3. It can be observed from Fig. 3(a) and (b) that each microsphere is composed of numerous primary particles with sizes of several tens of nanometers. In addition, HRTEM images from three different positions (regions (1), (2), and (3)) are displayed in Fig. 3(c) and (d). The apparent lattice fringes in region (2) with an interplanar spacing of 0.47 nm are observed and these fringes are assigned to the planar distance between the (003) planes of the layered LiCoO₂ and the clear lattice fringes in region (3) with a d-spacing of 0.203 nm correspond to the (104) plane of the α-NaFeO₂ layered structure.^26,27 Meanwhile, as for the introduction of Na ions into the Li slab, it can be seen from Fig. 3(c) and (d) that the Na-LMO sample exhibits a cubic shape with distinct edges and corners, which is possibly because introducing Na ions into the Li slab improves the stability of the layer structure so that the growth of the crystal tends to follow the direction of the layer. The lattice fringes with an interplanar spacing of about 0.47 nm are slightly widened, which leads to a more obvious layered structure as can be seen from region (1). A similar result of the widened crystal lattice brought about by the introduction of Na ions is subsequently observed from the XRD results, which is attributed to the introduction of Na ions into the alkali ion layer in the lattice, which substantially enlarges the Li slab space.

Fig. 3 (a) and (b) TEM images of the Na-LMO sample, (c) and (d) HRTEM images of the Na-LMO sample and the corresponding fast Fourier transform (FFT) images.

The crystal structures of the LMO and Na-LMO samples were identified with XRD patterns in Fig. 4. It can be clearly observed that all the peaks of the two samples can be indexed based on a hexagonal α-NaFeO₂ structure with space group R3̄m^6,7 except the weak superlattice peaks between 20° and 23°, which correspond to the C/2m derived from the short-range superlattice ordering of the Li, Ni, and Mn atoms in the transition metal layers.²⁶ Meanwhile, the two pairs of the (006)/(102) and (018)/(110) peaks of the two samples are well spilt, suggesting a well-defined layered structure formed in the lattice.^26–28 Based on a magnifying observation of the (003) peak in Fig. 4 (right), it can be seen that the (003) peak of the Na-LMO sample shows a slight shift toward the lower angle region. This phenomenon indicates that the Na ions have been completely introduced into the alkali ion layer in the lattice, which substantially enlarges the Li slab space.¹⁰

Fig. 4 XRD patterns for the LMO and Na-LMO samples.

In order to further characterize the specific surface areas and porous nature of the Na-LMO and R-LMO samples, nitrogen sorption isotherms were used to determine the pore volumes and PSD in terms of the BET equation and the BJH method. The N₂ adsorption–desorption isotherms and PSDs are shown in Fig. 5. It can be seen from Fig. 5 that the Na-LMO sample exhibits a type II with a type of H3 hysteresis loop, which is characteristic of microporous materials.^33,34 In addition, as can be seen from the corresponding BJH plots (the inset) recorded from the nitrogen isotherms of the Na-LMO sample, the average pore size is about 19.55 nm, confirming that the sample contains mesopores. However, it can be observed that the pore volumes of the R-LMO sample are much smaller than those of the Na-LMO, indicating that there are few pores in the R-LMO sample. It can be considered that the introduction of Na ions has a positive role in promoting the formation of the porous structure in these cathode microspheres, which is consistent with the results observed by the field emission SEM (FESEM)/TEM images (in Fig. 3).

Fig. 5 Nitrogen adsorption/desorption isotherms and pore size distribution curves (insets) of the (a) R-LMO and (b) Na-LMO samples.

In order to estimate the Na concentration in the composite and to determine the chemical states of the constituent elements, other characterization techniques, such as XPS, are still needed; the results of the two samples are exhibited in Fig. 6. All of the binding energies (BEs) in the XPS analysis are corrected for specimen charging by referencing them to the C 1s peak (set at 284.6 eV). It can be seen from Fig. 6 that the Ni 2p_3/2 binding energies of the two samples are in accordance with the standard Ni²⁺ and Ni³⁺, which are theoretically located at 854.1 eV and 855.1 eV, respectively.²⁶ In addition, the observed binding energies for Co 2p_3/2 (780 eV) and Mn 2p_3/2 (642.2) in the two samples coincide well with those for the standard Co³⁺ (779.9 eV) and Mn⁴⁺ (642.2 eV), respectively.³⁵ Besides, as can be seen from Fig. 6(b) and (f), the peak of Na 1s can be observed in the sample of Na-LMO. The peak position appears at a binding energy of approximately 1072.19 eV, which is the peak position of the binding energy of Na⁺ oxide, indicating that the Na element doped as Na⁺ in the sample of Na-LMO.^26,33 The XPS results indicate again that the Na element has been successfully introduced into the porous cathode microspheres as Na⁺, and the introduction of Na⁺ has no influence on the oxidation state of the transition metal ion.

Fig. 6 XPS spectra of (a) and (b) survey spectrum, (c) Mn 2p, (d) Ni 2p, (e) Co 2p, and (f) Na 1s for the Na-LMO and LMO samples.

As discussed above, a unique porous Na-doped microsphere Li-rich cathode material can be successfully obtained by a solvothermal method and subsequent high temperature solid state method. It will be interesting to study the electrochemical properties of the Na-LMO sample in lithium-ion batteries. The electrochemical performance of lithium cells with a porous Na-LMO electrode and LMO electrode are shown in Fig. 7. Fig. 7(a) and (b) show the charge/discharge curves of different cycles at 0.1C. It can be found that the initial discharge capacity of the Na-doping (Na-LMO) sample at 0.1C is about 305.3 mA h g⁻¹ with a high coulombic efficiency of 88.15% and a low irreversible capacity loss of 41.04 mA h g⁻¹, while the irreversible capacity loss for the R-LMO sample is 66.89 mA h g⁻¹ and the coulombic efficiency is only 81.2%. Similar to most Li-rich cathodes, all the cells exhibited two voltage plateaus around 4.0 and 4.5 V, corresponding to the extraction of lithium ions from the layered structure and the extraction of lithium and oxygen from the Li₂MnO₃ structure during the first charging process, respectively.^14,15 In the meantime, it can be obviously seen from Fig. 7(c) that the Na-LMO sample shows an excellent cycle stability compared with the LMO sample and a stable reversible discharge capacity of 268.9 mA h g⁻¹ and a capacity retention of 87.9% can be observed after 200 charge/discharge cycles. The capacity retention of the Na-LMO sample even reaches 93.1% after 100 cycles, which is apparently higher than the previously reported materials.^26,27 These improvements in the Na-LMO sample are ascribed to the more freedom of spaces for the transfer of lithium ions due to the existence of a “porous” structure and the much larger units created by the doping of Na ions.^10,15 Thus, the existence of the porous structure and the Na ion introduction lead to a better electronic pathway and a high rate capacity of the electrode active material. The rate performance of the lithium cells in the voltage window 2.0–4.6 V is shown in Fig. 7(d) and the cells are initially cycled between 2.0 and 4.6 V at 0.1C, 0.2C, 0.5C, 1C, 2C, 5C, and 10C for 6 cycles per step. Significantly, the porous Na-LMO electrode delivers a high discharge capacity of 177.2 mA h g⁻¹ even at 10C, whereas only 122.8 mA h g⁻¹ specific capacities can be obtained for the LMO electrode at the same rate. For the Na-LMO sample, porous materials with subunit nano-particles are important for the rate capacity and cyclic stability of the lithium ion battery. The nano-size particles reduce the path of lithium-ion diffusion and increase the number of reaction sites for lithium intercalation/deintercalation, and the interior pores provide spaces to buffer the volume changes during charge–discharge and thus enhance the rate capacity and cyclic stability of the materials.^15,16

Fig. 7 The charge/discharge curves of the (a) LMO electrode and (b) Na-LMO electrode at different cycles, (c) cycling performance of the LMO and Na-LMO electrodes, (d) rate performance of the LMO and Na-LMO electrodes.

In order to intensively explore the cause of the improved cycling ability of the Na-doped porous Li-rich cathode materials, electrochemical impedance spectroscopy (EIS) measurements were carried out for the Na-LMO and LMO samples. The corresponding results are shown in Fig. 8 and the corresponding data can be seen in Table 2. As can be seen, both of the two impedance spectra consist of two semicircles in the high and intermediate-frequency ranges and a line inclined at a constant angle to the real axis in the low-frequency range. According to previous EIS studies on this type of layered cathode material,^26,27 the first semicircle in the high-frequency region can be mainly ascribed to the Li⁺ ion diffusion through the SEI film, the second semicircle in the intermediate frequency range can be ascribed to the charge transfer reaction at the interface of the cathode–electrolyte, and the inclined slope in the low frequency region can be assigned to Li⁺ diffusion (Warburg impedance) in the electrode bulk.²⁵ On the basis of this understanding, R_s represents the internal resistance of the battery, R_f represents the resistance of the SEI film, R_ct represents the charge-transfer resistance, and W is the Warburg impedance related to Li⁺ diffusion.^25,26,30 It can be found from the inset in Fig. 8 and Table 2 that the semicircle in the high frequency district for the EIS plot of Na-LMO sample is closer to the Z_im axis than the LMO sample, and thus the value of R_s is less than the LMO sample, which signifies a decreased ohmic resistance (R_s) in the Na-LMO sample. In addition, as Na ions create larger unit cells (which has been discussed in previous XRD analysis), it causes better transition metal ion motion and/or migration, thereby improving conductivity through the correlated motion of electrons, thus, the sample with 5 at% doping (Na-LMO) can reciprocate a high conductivity.^28,31 It can be seen from Fig. 8 and Table 2 that the charge transfer resistance (R_ct) of the Na-LMO sample is much lower than that of the LMO sample. The Na-LMO sample also shows a lower charge transfer resistance after the 100^th cycle compared with the LMO sample, which indicates that the Na-LMO sample can effectively suppress the decomposition of the electrolyte to form a thinner SEI film and enhance the kinetics of lithium-ion diffusion due to the enlarged Li slab space. It is believed that the enlargement of the Li slab space by Na doping can facilitate the diffusion and removal of the lithium ion and transition metal ion in the bulk of the material.^10,32 In addition, the porous structure in the microspheres provides shorter diffusion pathways for ions as well as faster electron transport during cycling, which is attributed to the extra spaces within the microspheres.

Fig. 8 Electrochemical impedance spectra after (a) the 1^st cycle and (b) the 100^th cycle for the Na-LMO and LMO samples.

Table 2 The relative amounts of Li, Ni, Co, and Mn in Na-LMO and LMO

Sample	R-LMO		Na-LMO
	1st	100th	1st	100th
Rs	5.32	12.57	3.77	5.55
Rct	107.30	258.80	67.08	248.5
CPE	5.84 × 10^−6	5.35 × 10^−6	1.80 × 10^−5	8 × 10^−6
Cp	0.71	0.77	0.71	0.86

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4. Conclusions

In summary, 5 at% Na-doped porous Li-rich microspheres for advanced lithium-ion battery cathode material were prepared through the solvothermal method and a subsequent high-temperature solid state method. The cathode microspheres show a porous internal structure and 5 at% Na ions are doped. The introduction of Na ions and the porous structure in this new Li-rich cathode material can jointly improve its electrochemical behavior, resulting in a lower charge transfer resistance of the materials, which is associated with a higher electronic conductivity of the cathode. In addition, the Na-doped porous cathode material shows a high initial discharge capacity of 305.3 mA h g⁻¹ with a high coulombic efficiency of 88.2%. It even delivers a high discharge capacity of 268.9 mA h g⁻¹ after 200 cycles with a high capacity retention of 87.9%. Therefore, the introduction of Na ions into the Li layer can achieve an enhanced lithium ion diffusion and improved stability of the layered structure, and the existence of the porous structure can provide extra active sites for Li⁺ storage, and improvements in the electrochemical performance can be realized based on these modifications.

Acknowledgements

This work is supported financially by the National Natural Science Foundation of China under project No. 51472211, Scientific and Technical Achievement Transformation Fund of Hunan Province under project No. 2012CK1006, Key Project of Strategic New Industry of Hunan Province under project No. 2013GK4068, and the Natural Science Foundation of Hunan Province under project No. 2015JJ6103. National Science Foundation for Post-doctoral Scientists of China under project No. 2015M570682. Scientific Research Foundation of the Higher Education institutions of Hunan Province under project No. 15C1313.

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