A further electrochemical investigation on solutions to high energetical power sources: isomerous compound 0.75Li_1.2Ni_0.2Mn_0.6O₂·0.25LiNi_0.5Mn_1.5O₄

Abstract

An isomerous layered/spinel 0.75Li_1.2Ni_0.2Mn_0.6O₂·0.25LiNi_0.5Mn_1.5O₄ cathode material with outstanding electrochemical properties has been synthesized by a reasonable design of introducing high-power spinel LiNi_0.5Mn_1.5O₄ material to fill up the surface gaps of pristine lithium-rich layered Li_1.2Ni_0.2Mn_0.6O₂ material with a molar ratio of 25 : 75. Morphological characterization reveals that the octahedral spinel LiNi_0.5Mn_1.5O₄ particles are successfully coated into the surface gaps of the Li_1.2Ni_0.2Mn_0.6O₂ secondary particle, forming a special alternant structure with spherical and octahedral particles on the surface. Interestingly, some hollow sections are also observed in 0.75Li_1.2Ni_0.2Mn_0.6O₂·0.25LiNi_0.5Mn_1.5O₄ material, confirmed from the TEM images. The structural characterization demonstrates that this isomerous compound is more well-defined α-NaFeO₂ configured, more enlarged in Li layer spacing and lower in cation disordered degree. The exquisite morphology and ideal structure endow this nanocrystal-assembled composite significantly enhanced electrochemical performance with high capacity, good rate capability and excellent cycling stability, compared with the pristine Li_1.2Ni_0.2Mn_0.6O₂. It delivers a discharge capacity of 135 mA h g⁻¹ even at an ultrahigh current density of 2000 mA g⁻¹ (10 C). Moreover, the superior cycling stability is also observed with high discharge capacities of 254 mA h g⁻¹ and 222 mA h g⁻¹ at 0.5 C and 1 C after 100 cycles with capacity retention of 98% and 94%, respectively. Moreover, the fast-charging test results are indicative of the fact that this layered/spinel cathode could be used in practical application. Its discharge capacity is 176 mA h g⁻¹ at 1 C after 50 cycles with the charge rate of 10 C. Furthermore, the composite can endure high current charging and discharging even at a high cut-off potential (5.0 V), whereas the pristine Li_1.2Ni_0.2Mn_0.6O₂ material cannot. Therefore, we absolutely believe that this isomerous layered/spinel 0.75Li_1.2Ni_0.2Mn_0.6O₂·0.25LiNi_0.5Mn_1.5O₄ cathode is a promising candidate for the commercial development of advanced LIBs.

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1. Introduction

To meet the increasingly demanding requirements for the storage and utilization of clean energy, lithium-ion batteries are now being intensively developed as alternative power sources for hybrid electric vehicles (HEVs) and electric vehicles (EVs) due to their high energy- and high power-density and environmental benignancy.^1–7 Layered lithium-rich materials Li[M_1−xLi_x]O₂ (M = Mn, Ni, Co), which are solid solutions between the monoclinic Li₂MnO₃ (C2/m symmetry) phase and trigonal LiMO₂ (R3̄m symmetry) phase,^8–11 have attracted particular attention because of their low cost and encouraging high capacity (∼250 mA h g⁻¹) when charged above 4.5 V.^12–14 Despite these advantages, their intrinsically inferior rate capability, poor cycle performance and the initial large irreversible capacity loss (ICL), accompanied with the dissolution of transition metal ions, erosion from the electrolytes and fragile surface properties at high potential, have prevented them from practical applications.^15–18

Accordingly, many efforts, such as acid treatment, using ionic liquid electrolyte, nano-architecture fabrication and surface modification, have been made to handle these shortcomings.^19–25 For example, S.-H. Kang et al. reported that the acid treatment of the lithium-rich electrodes can eliminate the initial irreversible capacity loss but worsen their cycling stability and rate capability.¹⁹ Surface coating with AlF₃ was reported to promote the cycling performance as well as the thermal stability due to the spinel structure, which was partially introduced by the AlF₃ coating, can suppress the increase of charge transfer resistance (R_ct).^23,24 Moreover, coating with Li⁺-conducting phosphate LiNiPO₄ shows excellent rate capability (200 mA h g⁻¹ at 1 C).²⁵ However, these coating materials are electrochemically inactive, thus leading to a limited improvement of electrochemical performances. Hence, it remains a big challenge to apply these layered lithium-rich cathodes for high energy density batteries.

Currently, integrating the high power spinel material with the high capacity lithium-rich layered material to fabricate a layered/spinel composite is considered to be a feasible method to obtain the high-power-density cathode material. The spinel oxides possess high Li⁺ diffusion coefficients and 3D Li⁺ diffusion channels, which can facilitate fast Li⁺ transport between the bulk material and the electrolyte, so the rate performance is excellent. However, the spinel cathode delivers a capacity of less than 150 mA h g⁻¹ at above 3 V cycling. On the other hand, the lithium-rich layered oxides can exhibit a high discharge capacity at low rate. Furthermore, the cubic close-packed oxygen arrays in both the spinel and the layered oxides are structurally compatible.^17,26 Therefore, combining the merits of the layered and the spinel materials to render a layered/spinel composite could be an effective way to simultaneously achieve high capacity and outstanding rate capability as well as enhanced cycle performance.

Based on these considerations mentioned above, we develop a distinct design and synthesis of a layered/spinel isomerous 0.75Li_1.2Ni_0.2Mn_0.6O₂·0.25LiNi_0.5Mn_1.5O₄ material, using a high-energy spinel LiNi_0.5Mn_1.5O₄ material to fill up the surface gaps of pristine lithium-rich layered Li_1.2Ni_0.2Mn_0.6O₂ material. Results reveal that the obtained 0.75Li_1.2Ni_0.2Mn_0.6O₂·0.25LiNi_0.5Mn_1.5O₄ material is the spherical aggregates composed of spherical layered Li_1.2Ni_0.2Mn_0.6O₂ particles and octahedral spinel LiNi_0.5Mn_1.5O₄ particles closely arranged alternately on these secondary particles' surfaces. Moreover, the TEM images depict that some amazing hollow sections exist in this composite. This graceful morphology can maximize the inherent advantages of the 3D Li⁺ insertion/extraction framework of the spinel structure and high Li⁺ storage capacity of the layered structure. Encouragingly, this layered/spinel electrode shows excellent cycle stability and high-rate discharging performance as well as outstanding fast-charging ability. More surprisingly, this composite can also charge–discharge at high-rate between 2.0 and 5.0 V.

2. Experimental section

2.1 Sample preparation

The pristine layered lithium-rich cathode Li_1.2Ni_0.2Mn_0.6O₂ (PL) was prepared by a solid-state reaction from the Ni_0.25Mn_0.75CO₃ precursor and Li₂CO₃. The Ni_0.25Mn_0.75CO₃ precursor was prepared by co-precipitation method, and the corresponding synthetic process was reported in our previous work.²⁷ A 1.5 M mixture of NiSO₄·6H₂O and MnSO₄·H₂O (Ni : Mn = 1 : 3, molar ratio) was dissolved in distilled water and then pumped into a continuous stirred tank reactor (CSTR). Simultaneously, the 3 M Na₂CO₃ and 4.5 M NH₃·H₂O, as the precipitant and chelating agent, respectively, were separately fed into the CSTR at 50 °C. The pH value was carefully controlled at 8.0 ± 0.2, and the stirring speed was monitored at around 1000 rpm. Co-precipitated powders were filtrated and washed multiple times and finally dried at 100 °C for overnight. Thereafter, the obtained precipitates were mixed with 6% excess Li₂CO₃, and first preheated at 550 °C for 6 h and then calcined at 900 °C for 10 h in air.

To obtain the isomerous 0.75Li_1.2Ni_0.2Mn_0.6O₂·0.25LiNi_0.5Mn_1.5O₄ (SL) cathode, Ni(NO₃)₂, Mn(CA)₂, C₁₂H₂₂O₁₁ and LiNO₃ (excess 5%) were first dissolved in distilled water. Then, the as-synthesized PL particles were dispersed in this solution by CSTR. This mixture solution was stirred vigorously at 80 °C for 12 h to form homogeneous dispersoid and then dried in the vacuum oven at 60 °C. Afterwards, the dried powders were calcined at 900 °C for 6 h, and then re-annealed at 700 °C for 6 h in air. The preparation route is illustrated in Scheme 1.

Scheme 1 Schematic of the preparation route of 0.75Li_1.2Ni_0.2Mn_0.6O₂·0.25LiNi_0.5Mn_1.5O₄ material.

2.2 Sample characterization

The structure of the as-prepared compounds was characterized by X-ray diffraction (XRD, PW1730, Cu Kα, λ = 1.5418 Å). The XRD data were collected in the range of 15–70° (2θ). Structural refinement was performed using TOPAS software. Particle morphologies and size of the as-prepared samples were measured by scanning electron microscopy (SEM, S4800). The valence states of manganese and nickel in the samples were examined by X-ray photoelectron spectra (XPS, ESCALAB 250Xi, Al-Kα). The microstructures of these composites were observed by transmission electron micrograph (TEM, JEM-2100).

2.3 Electrochemical measurement

Electrochemical measurements were made using CR2032-type coin cells consisting of the as-prepared oxide cathode, Li metal anode and a porous polypropylene separator (Celgard 2400) with 1 M LiPF₆ solution in EC-DMC (1 : 1 in volume rate) as the electrolyte. The cathodes were prepared by mixing oxide powder, carbon black and polyvinylidene difluoride (PVDF) binder (80 : 13 : 7, weight ration) in N-methylpyrrolidone (NMP). The obtained slurry was casted on an aluminum foil and dried in the vacuum oven at 100 °C for 10 h. The dried cathode sheet was cut into discs with diameters of 14 mm and then pressed under a pressure of 20 MPa. The loading of the active material in the electrode was 3–4 mg cm⁻². All coin cells were assembled and sealed in an argon-filled glove box. The cells were charged and discharged galvanostatically at room temperature with different current densities (1 C = 200 mA g⁻¹) in the voltage range of 2.0–4.8 V and 2.0–5.0 V (vs. Li/Li⁺). Electrochemical impedance spectroscopy (EIS) measurements were tested using an electrochemical workstation (Zennium, IM6) in the frequency range of 100 kHz–10 mHz with an alternating-current amplitude of 5 mV.

3. Results and discussion

3.1 Morphology and structure characterizations

Fig. 1 shows the Rietveld refinement results for XRD patterns of the PL and SL samples. As shown in Fig. 1a, the strong peaks of the PL sample can be indexed by α-NaFeO₂ rock structure with space group R3̄m, which is normally taken as the layered characteristic of LiMO₂ (M = Ni, Mn) phase.^8,28 In addition, a series of weak reflections at 21–25° (2θ value) are in accordance with the hexagonal LiMn₆ super-ordering in Li₂MnO₃ monoclinic phase with C2/m symmetry.^8,28 To distinguish the two phases in graphs, they are marked by “R” and “C”. Fig. 1b shows four phases in the SL sample: LiMO₂ layered phase, Li₂MnO₃ monoclinic phase, spinel disordered phase (Fd3̄m) and ordered phase (P4₃32) of LiNi_0.5Mn_1.5O₄. The Fd3̄m phase in spinel LiNi_0.5Mn_1.5O₄ material can be reasonably attributed to the insufficient oxygen synthesis condition, resulting in some partial product incomplete oxidation.^29,30 The two different spinel phases are also marked by “F” and “P” in Fig. 1b.

Fig. 1 Rietveld refinement of the X-ray diffraction patterns for (a) PL and (b) SL samples.

The refined lattice parameters of PL and SL samples are given in Table 1. The lattice constant a is slightly decreased from 2.8604 Å (PL) to 2.8484 Å (SL). However, the lattice constant c is increased from 14.2505 Å (PL) to 14.2656 Å (SL). The c-lattice parameter is perpendicular to the Li layer in the layered structure. Therefore, its increase signifies an enlargement of Li layer spacing.³¹ In addition, the Li slab space can be defined as the average distance between the oxygen layers around the Li layer. The increased Li slab space could cause a substantially higher Li diffusivity.³¹ The c/a values of PL and SL are 4.9820 and 5.0083, respectively, which are both higher than 4.899 for ideal hexagonal α-NaFeO₂ structure,^32,33 indicating a more well-defined layer structure for the SL sample. The cation disordering of Li⁺ and Ni²⁺ can be estimated by the I₍₀₀₃₎/I₍₁₀₄₎ ratio: the greater it is, the less the disordering is.^32–34 The I₍₀₀₃₎/I₍₁₀₄₎ ratio of SL (1.88) is greater than that of PL (1.72), indicating a lower level of cation disordering for SL. The weight ratio of different phases in the PL and SL samples are also shown in Table 1. The ratio of layered LiMO₂ (R3̄m) and layered Li₂MnO₃ (C2/m) in PL is very close to 1 : 1. This result is in good agreement with the previous reports about this cathode material.^35–37 Similarly, the weight ratio of layered LiMO₂ (R3̄m) and layered Li₂MnO₃ (C2/m) in SL is also approximately 1 : 1, suggesting that the pristine lithium-rich layered structure is not broken in the SL sample. In addition, the weight ratio of the layered phase (R3̄m and C2/m) and spinel phase (Fd3̄m and P4₃32) in the SL sample is 74.79 : 25.21, which approaches the expected ratio and demonstrates the successful synthesis of 0.75Li_1.2Ni_0.2Mn_0.6O₂·0.25LiNi_0.5Mn_1.5O₄ composite.

Table 1 Refined lattice parameters of PL and SL samples

Sample	a (Å)	c (Å)	c/a	I(003)/I(104)	Weight (%)				Rwp (%)	Rp (%)
					Layered (R3̄m)	Layered (C2/m)	Spinel (Fd3̄m)	Spinel (P4332)
PL	2.8604	14.2505	4.9820	1.72	49.894	50.106	—	—	2.18	1.70
SL	2.8484	14.2656	5.0083	1.88	37.768	37.019	15.504	9.709	2.39	1.87

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X-ray photoelectron spectroscopy (XPS) measurements are employed to elucidate the variation in chemical states. Fig. 2 shows the typical XPS spectra of Ni 2p and Mn 2p for the PL and SL samples. All XPS spectra are corrected using C 1s at 284.60 eV. In the PL sample, the observed binding energies at 642.2 eV and 854.5 eV are consistent with the previous report for Mn⁴⁺ and Ni²⁺ in similar oxide cathode materials, respectively.^38–40 However, the Mn 2p_3/2 and Mn 2p_1/2 peaks are slightly shifted toward lower binding energies for the SL sample, indicating some lower oxidation state than tetravalent Mn ions that appeared. The variation of Mn oxidation in the SL sample can be reasonably ascribed to the spinel LiNi_0.5Mn_1.5O₄ (Fd3̄m) phase that exists, in which a small amount of manganese remains as Mn³⁺ in consequence of the incomplete oxidation of manganese during the high temperature calcination.^29,30 Correspondingly, a similar peak shift of Ni 2p spectra toward higher binding energies is also observed in the SL sample. It is probably that the increase in oxidation state of Ni is a compensation of deduction in oxidation state of Mn ions.

Fig. 2 XPS of Ni 2p and Mn 2p spectra for PL and SL samples.

Scanning electron microscopy (SEM) images of the as-prepared PL and SL materials are shown in Fig. 3. As seen from these pictures, there are significant differences in the morphologies of the two samples. The PL sample (Fig. 3a–d) consists of homogeneously distributed secondary spherical particles with an average size of about 2 μm (Fig. 3b). Each secondary particle is composed of sphere-like particles with a small size of about 80 nm, as shown in Fig. 3c. Interestingly, this second spherical morphology is preserved in the SL sample (Fig. 3e–h). However, from comparison, some remarkable changes can be observed: (i) the secondary particles are assembled by primary octahedral shape particle and spherical particle, whereas the PL is made only by the primary spherical particle; (ii) the primary particles are increased to 140 nm (Fig. 3g), and the average size of secondary particles are increased to 3 μm (Fig. 3f); (iii) the secondary particles become more dense due to the octahedron spinel particles occupying the surface porosity of the PL sample after the LiNi_0.5Mn_1.5O₄ phase is introduced. In addition, the transmission electron micrograph (TEM) images of the single secondary particle of PL and SL are shown in Fig. 3d and h, respectively. These two samples present the same sphere morphology, except that SL has some octahedron-shaped particles on the surface, marked by red circles. Particularly, some bright cavities are observed in both TEM images, marked by yellow arrows, indicating that some hollow sections exist in the particle center. Some studies have confirmed that this microsphere with hollow interior can boost the rate capability and cycling stability significantly.^41–43 Because this type of structure can provide a shorter path for Li⁺ ion diffusion and a larger electrode–electrolyte contract area for Li⁺ flux across the interface, the structural integrity can be maintained by alleviating the mechanical strain that occurs during the repeated Li⁺ ion insertion/extraction processes.

Fig. 3 SEM and TEM images of (a–d) PL sample and (e–h) SL sample.

Fig. 4 shows the high resolution transmission electron micrograph (HRTEM) and the selected area electron diffraction (SAED) of as-synthesized PL and SL materials. The HRTEM of the PL sample (Fig. 4a) shows a typical layered structure with lattice fringes spacing of 0.47 nm, which matches well to the (003)_Hex plane of LiMO₂ (M = transition metals) and/or (001)_Mon plane of Li₂MnO₃. Correspondingly, the SAED along the [2̄021]_Hex zone axis consists of two sets of reflections: hexagon that is formed by six neighboring bright spots and some weak spots in the hexagon. As reported,^28,44 the weak spots manifest the presence of a Li₂MnO₃-like phase and the ordering of Li ions with TM ions in TM layers. The pattern of spots marked by yellow circles is for the layered LiMO₂ phase, and the pattern of spots connected by red lines is for the monoclinic Li₂MnO₃ phase. As seen from the SEM and TEM images of the SL sample, shown in Fig. 3, we can reasonably deduce that this material is mainly made up by sphere particles and a handful of octahedral-shape particles. Therefore, the two different morphologies particles are selected to investigate the nano-structure. Fig. 4b shows the HRTEM image and SAED patterns of the primary octahedral-shape particle. The lattice fringes are measured to be 0.47 nm, which is well indexed to the (111)_Cub plane of the spinel LiNi_0.5Mn_1.5O₄ phase.^30,45 In addition, the SAED pattern is consistent with a typical spinel lattice structure, which is collected from a crystal analyzed along the [101̄]_Cub zone axis. Similarly, the HRTEM and SAED patterns of the primary sphere particle are shown in Fig. 4c. The lattice fringe has a distance equal to 0.42 nm, which coincides with the interplanar distances of the (020)_Mon plane of Li₂MnO₃ phase. The SAED pattern exhibits strong hexagonal reflections, indexed in six-index notation (marked by yellow circle), indicating a α-NaFeO₂ layered structure (R3̄m).⁴⁶ The weak reflections (marked by red circles), which are associated with Li ordering in Li₂MnO₃-like domains, are clearly identified between the two bright spots of the layered trigonal symmetry (R3̄m). This SAED result is consistent with the analysis of the PL sample (Fig. 4a), indicating that the sphere particle of the SL sample is also a lithium-rich layered structure. Therefore, it is reasonable to conclude that the pristine lithium-rich layered structure is not damaged in the SL sample by introducing the spinel LiNi_0.5Mn_1.5O₄ phase.

Fig. 4 HRTEM and SAED patterns of (a) PL and (b), (c) SL samples.

3.2 Electrochemical performance

The initial charge–discharge profiles of PL and SL samples in the voltage window of 2.0–4.8 V are shown in Fig. 5a. By comparison, it could be found that both electrodes deliver the similar charge capacity of about 350 mA h g⁻¹ at 0.1 C. Moreover, the PL sample suffers from large irreversible capacity loss, which is attributed to the irreversible removal of Li₂O during Li₂MnO₃ activation above 4.5 V.^12–14 The SL sample exhibits relatively lower irreversible capacity loss than the PL sample. As shown in Table 2, the ICL is 54 mA h g⁻¹ for SL, compared with 81 mA h g⁻¹ for PL. Combined with the anomalous observation of two newly formed discharge plateaus at 4.7 V and 2.7 V (marked by the yellow circle) in the SL sample, which are frequently recognized as the characteristics of the spinel phase,⁴⁷ the increased discharge capacity as well as the enhanced first coulomb efficiency could be ascribed to the unoccupied 16c sites of spinel LiNi_0.5Mn_1.5O₄ phase, which can accommodate the additional Li⁺ ions that are lost as a consequence of the activation of Li₂MnO₃ during the initial charge process of layered lithium-rich component.²⁶ To clearly investigate the impact of spinel LiNi_0.5Mn_1.5O₄ phase on the electrochemical reactions, the corresponding dQ/dV plots are presented in the inset. The PL sample presents a typical lithium-rich layered dQ/dV curve.⁴⁸ The anodic peak around 4.5 V corresponds to Li₂MnO₃ activation, causing irreversible oxidation of O_2p from the cathode (removed as a formula of Li₂O).⁴⁹ Another anodic peak is observed around 3.8 V corresponding to the oxidation of Ni²⁺ to Ni⁴⁺. For the SL sample, however, the sharp peak at around 4.5 V has a lower intensity than the PL sample due to the intrinsic low content of the Li₂MnO₃ component. Moreover, more surprisingly, the Li₂MnO₃ anodic peak is shifted toward high voltage after introducing high voltage spinel LiNi_0.5Mn_1.5O₄ material (marked by aqua green dotted line). The anodic peak of layered Ni^2+/4+ is replaced by the spinel Ni^2+/4+ peak at around 4.7 V in the SL sample, and the oxidation of Ni²⁺ to Ni⁴⁺ is split into two continuous reactions of Ni²⁺ → Ni³⁺ and Ni³⁺ → Ni⁴⁺, respectively. In addition, a small anodic peak at about 2.9 V is also observed, which can be ascribed to the inter-conversion between the first-order spinel LiNi_0.5Mn_1.5O₄ phase and the rock salt Li₂[Ni_0.5Mn_1.5]O₄ phase.⁵⁰

Fig. 5 (a) Initial charge–discharge profiles and the corresponding charge dQ/dV plots (inset); (b) capacity contribution of first cycle in different voltage range of PL and SL samples.

Table 2 Initial charge–discharge electrochemical data of PL and SL samples

Sample	Charge capacity (mA h g^−1)	Discharge capacity (mA h g^−1)	Irreversible capacity loss (mA h g^−1)	Coulomb efficiency (%)
PL	343	262	81	76
SL	347	293	54	84

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Capacity contributions of the first cycle in different voltage ranges for PL and SL samples are shown in Fig. 5b. The corresponding values of capacity contribution are summarized in Table S1 in ESI.† The PL and SL samples are capable of providing 219 mA h g⁻¹ and 240 mA h g⁻¹ charge capacities above 4.4 V with 64% and 69% contributions to the total charging capacity. Clearly shown in Fig. 5a, the intensity of the Li₂MnO₃ anodic peak of PL at around 4.5 V is higher than SL, indicating the more capacity contribution of Li₂MnO₃ component for PL. However, the charge capacity derived from above 4.4 V in SL is still stronger than PL; therefore, we can reasonably deduce that the extra-charge capacity is attributed to the oxidation of Ni²⁺ to Ni⁴⁺ around 4.7 V for the spinel phase. The discharge capacity contribution from lower than 3.5 V part in SL is 162 mA h g⁻¹, which is higher than 125 mA h g⁻¹ in PL. As mentioned in previous reports, this result indicates more active redox of Mn^4+/(4−x)+ reactions in SL.¹⁸

To the best of our knowledge, the poor rate capability and unstable cyclic performance of the layered lithium-rich material are the main obstacles to their large-scale applications in electric vehicles (EVs). However, the following facts would prove that our work has prepared a layered/spinel composite (SL) with outstanding high-rate discharge capacity and cycle performance. Fig. 6 shows the rate capability of PL and SL samples with the incremental rates from 0.1 C to 10 C between 2.0 and 4.8 V. Evidently, the SL sample exhibits more excellent rate performance. It yields high maximal discharge capacities of 293, 279, 261, and 238 mA h g⁻¹ at 0.1 C, 0.2 C, 0.5 C and 1 C, respectively, and more surprisingly, a maximal capacity of 135 mA h g⁻¹ is still achieved at 10 C. Such superior rate capability might be related to the inherent advantage of the higher Li⁺ conductivity of the spinel LiNi_0.5Mn_1.5O₄ phase. This improved electrochemical performance is comparable to or better than some previous excellent reports on this type of layered/spinel composite,^26,51 as demonstrated by a recent case in which a layered/spinel composite Li_1.3Ni_0.25Mn_0.75O_2.4 (0.1LiNi_0.5Mn_1.5O₄·0.8(Li₂MnO₃·LiNi_0.5Mn_0.5O₂)) shows an excellent rate capability (102 mA h g⁻¹ at 10 C) and cycling stability (98.2% for 40 cycles at 0.2 C) by introducing the spinel structure into the layered lithium-rich system.⁵¹ In contrast, the PL sample shows an inferior rate performance. It delivers a maximal discharge capacity of 271 mA h g⁻¹ at 0.1 C but fast fades to 82 mA h g⁻¹ at 10 C. The serious capacity fading results from fast Li⁺ ions intercalation/deintercalation, which destroys the fragile surface of the layered lithium-rich structure at high rates.^8,16 Furthermore, some of this deterioration of electrochemical behavior in the Li metal half cells could be associated with the significant changes on the anode Li side, which has been systematically investigated elsewhere.⁵²

Fig. 6 Rate capability of PL and SL samples between 2.0 and 4.8 V.

The long-term cycling tests of PL and SL electrodes after rate capacity examination are shown in Fig. 7. After 100 cycles, the PL sample shows a discharge capacity of 176 mA h g⁻¹ at 0.5 C with acceptable capacity retention of 85%. However, it only retains a maximal capacity of 148 mA h g⁻¹ at 1 C with a poor capacity retention of 78%. However, the SL sample exhibits excellent cycling stability with good capacity retention. The discharge capacity at 0.5 C approaches as high as 259 mA h g⁻¹ at first cycle, maintaining 254 mA h g⁻¹ after 100 cycles. Even at a larger rate of 1 C, it still retains 94% of the initial capacity (236 mA h g⁻¹) after 100 cycles. To confirm the effects of introducing the spinel structure on the voltage fading upon cycling, which is one of the key challenges for high-capacity lithium-rich composite cathode,^36,48 the charge–discharge curves of PL and SL electrodes at different cycle time during the 100 cycles (Fig. 7) are shown in Fig. S1 in the ESI.† From the comparison it can be clearly observed that the SL sample exhibits a slower decrease in the plateau of discharge curves during the long-term cycling. This analysis indicates that the inherent defect of voltage fading upon cycling can be alleviated in the SL sample by coating some spinel structure on the layered lithium-rich material.

Fig. 7 Cycling stability of (a) PL and (b) SL samples after rate capacity examination.

Apart from the rate performance and cycle stability, the fast-charging capability is also very important for practical applications. Therefore, tests based on 10 C charge and 1 C discharge after an initial 0.1 C charge and discharge cycle are carried out on both PL and SL, as shown in Fig. 8. Most surprisingly, a capacity of about 225 mA h g⁻¹ is obtained first by the SL sample and then gradually fades to 153 mA h g⁻¹ at the 100^th cycle. In contrast, the PL sample provides not more than 145 mA h g⁻¹ at first cycle and finally ends with 50 mA h g⁻¹ after 100 cycles. Such outstanding overall rate performance of SL material can be ascribed to the as-designed special structure, in which the spinel particles and layered particles are arranged alternately on the secondary particle surface. This distinctive surface structure can provide high Li⁺ conductivity and facilitate fast Li⁺ diffusion from electrolytes to the inter-layered structure by the three-dimensional (3D) diffusing channels in the spinel phase. Moreover, the compatibility and integration of the cubic spinel structure and layered structure on the surface ensure a stable structure to this layered/spinel material. In addition, the close-packing way on the surface can inhibit the internal layered lithium-rich material from erosion of electrolytes and restrain the bulk active-mass loss.

Fig. 8 10 C charge and 1 C discharge after an initial 0.1 C charge–discharge cycle between 2.0 and 4.8 V.

To investigate the capability of enduring high potential of this layered/spinel material, the discharge capacities of the two electrodes are tested at increasing rates from 0.1 C to 10 C and then recovering back to 0.2 C between 2.0 and 5.0 V. As can be seen in Fig. 9, the capacity of the PL sample is abruptly decreased with the increased current density and is finally broken at 3 C. Conversely, the SL sample is able to withstand the high potential, and it exhibits a maximal capacity of 78 mA h g⁻¹ at 10 C and subsequently returns to 0.2 C with the capacity of 181 mA h g⁻¹. This result indicates, once again, that the surface coating with spinel LiNi_0.5Mn_1.5O₄ phase endows this layered/spinel material with high structural stability. However, we can find that the discharge capacity of both samples is significantly decreased, compared with the situation of 4.8 V cut-off voltage. The reason can be partly attributed to the more significant decomposition of electrolyte at such a high cut-off voltage of 5 V.⁵³ Furthermore, the charge–discharge curves of both samples at various rates are shown in Fig. S2.† The image shows that the charge–discharge curve-shape changes greatly at different cycles and the plateau of discharge curves decreases with the increasing cycle number. Generally, this phenomenon is consistent with the bulk layered structure transformation into spinel phase during cycling and the increasing of inner impedances.^18,49

Fig. 9 Rate capability of PL and SL samples between 2.0 and 5.0 V.

Electrochemical impedance spectroscopy (EIS) has been performed to elucidate the difference in electrochemical performance of PL and SL samples. The measurements are carried out after rate capacity examination (Fig. 6). All Nyquist plots are shown in Fig. 10, and the corresponding equivalent circuit is presented in the inset. In this equivalent circuit, R_s represents the ohmic resistance, R_ct corresponds to charge transfer resistance, and W₁ is the Warburg impedance for depicting Li⁺ ion diffusion in the bulk structure.^18,48 By monitoring the diameter of the semicircles at high frequency, it is found that the R_ct for SL is smaller than that of PL, which indicates an improvement in the kinetics of Li⁺ diffusion through the surface layer and charge transfer reaction and a consequent increase in rate performance.⁵⁴ The considerably lower R_ct value of SL can be ascribed to three reasons: (i) the surface spinel phase LiNi_0.5Mn_1.5O₄, which possesses 3D channels for fast Li⁺ diffusion, whereas the pristine layered surface only has 2D channels; (ii) the surface spinels of SL can promote the Li⁺ ion conductivity and effectively reduce the barrier for Li⁺ ion transfer at the electrode–electrolyte interface; (iii) Li₂MnO₃ component is widely considered electrochemically inactive because Mn⁴⁺ ion is hardly oxidized in an octahedral coordination therefore leading to the lower content of Li₂MnO₃ in the SL compared to that of PL, which results in lower charge transfer reaction resistance (R_ct).⁵⁵

Fig. 10 EIS plots of the PL and SL electrodes with an equivalent circuit (inset).

4. Conclusions

An isomerous 0.75Li_1.2Ni_0.2Mn_0.6O₂·0.25LiNi_0.5Mn_1.5O₄ compound is successfully synthesized, in which the primary spherical particles (Li_1.2Ni_0.2Mn_0.6O₂) and octahedral particles (LiNi_0.5Mn_1.5O₄) arrange alternately on the secondary particles' surfaces. More interesting, some hollow sections are also observed in the center of the particle, indicating a superior structure for the penetration of electrolyte and the transmission of electron and lithium ion. The XRD Rietveld refinement results demonstrate that the 0.75Li_1.2Ni_0.2Mn_0.6O₂·0.25LiNi_0.5Mn_1.5O₄ composite has a more ordered α-NaFeO₂ structure, enlarged Li layer spacing and reduced degree of cation mixing. The exquisite morphology and ideal structure afford this isomerous cathode material's excellent electrochemical performance. It delivers a remarkable cycling stability with 98% and 94% of capacity retention at 0.5 C and 1 C after 100 cycles, respectively. Moreover, it exhibits a superior rate capability with a high discharge capacity of 135 mA h g⁻¹ even at 10 C. Apart from the enhanced rate performance and cycle stability, the fast-charging ability is also very outstanding for 0.75Li_1.2Ni_0.2Mn_0.6O₂·0.25LiNi_0.5Mn_1.5O₄, which provides a discharge capacity of 176 mA h g⁻¹ at 1 C after 50 cycles when using 10 C charge and 1 C discharge. Particularly, the layered/spinel composite can be used at high voltage (5.0 V) without ruining the host structure, which indicates high structural stability. This work suggests that this isomerous compound 0.75Li_1.2Ni_0.2Mn_0.6O₂·0.25LiNi_0.5Mn_1.5O₄ could be a potential candidate material for the application of rechargeable Li-ion batteries in the rapid development of hybrid electric vehicles (HEVs) and electric vehicles (EVs).

Acknowledgements

This work was supported by the Science and Technology Pillar Program of Sichuan Province (2014GZ0077), the Youth Foundation of Sichuan University (no. 2011SCU11081), the Doctoral Program of Higher Education of China (no. 20120181120103), and the Open Foundation of National Engineering Center for Phosphorus Chemical Industry (2013LF1012).

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Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c5ra03289j

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