Research Article Cite This: ACS Sustainable Chem. Eng. 2018, 6, 4625−4632
pubs.acs.org/journal/ascecg
Enhanced Rate Performance of Al-Doped Li-Rich Layered Cathode Material via Nucleation and Post-solvothermal Method Wenchao Yan,† Yu Xie,† Jicheng Jiang,†,‡ Deye Sun,† Xiaodi Ma,† Zhenggang Lan,*,† and Yongcheng Jin*,† †
Qingdao Industrial Energy Storage Technology Institute, Qingdao Institute of Bioenergy and Bioprocess Technology, Chinese Academy of Sciences, Qingdao 266101, P.R. China ‡ University of Chinese Academy of Sciences, 19A Yuquanlu Road, Beijing 100049, P.R. China S Supporting Information *
ABSTRACT: Al-doped layered cathode materials Li1.5−xAlxMn0.675Ni0.1675Co0.1675O2 have been successfully synthesized via a rapid nucleation and post-solvothermal method. The surface morphology and crystal structures of Aldoped Li-rich materials are investigated via scanning electron microscopy, X-ray diffraction, Raman spectra, and X-ray photoelectron spectroscopy. After optimization, the Li1.45Al0.05Mn0.675Ni0.1675Co0.1675O2 (Al = 0.05) sample showed excellent electrochemical performance, and the discharge capacities are 323.7 and 120 mAh g−1 at a rate of 0.1 and 20 C, respectively. These improvements, based on electrochemical performance evaluation and density functional theory calculations, might be ascribed to the increased electron conductivity of layered Li-rich material via Al3+ ions doped into a crystal structure. KEYWORDS: Lithium ion batteries, Li-rich material, Al doping, High rate performance, Density functional theory
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INTRODUCTION Lithium ion batteries (LIBs) have been extensively investigated and widely used as rechargeable power sources for applications in energy storage fields and electric vehicles.1−3 The continued rapid growth of electrical vehicles requires rechargeable power sources with higher energy and power density; however, current secondary lithium ion batteries cannot meet the requirement. It is reported that positive electrodes restricted the energy stored in LIBs.4 At present, the common positive materials mainly include layered LiCoO2 (140 mAh g−1),5 olivine LiFePO4 (170 mAh g−1),6 and spinel LiNi0.5Mn1.5O4 (146 mAh g−1),7 but these materials still suffer from low specific capacity, hindering their utilization as a potential material in future electric vehicles. Li-rich layered cathode materials, written as xLiMO2·(1−x)Li2MnO3, have been extensively studied owing to their high specific capacity (250 mAh g−1) with a higher operational window of 4.6−4.8 V.8,9 At the same time, high contents of Mn and Ni in these materials provide the advantage of low cost.10 The Li-rich layered material with eminent properties has been deemed to be one of the most promising next-generation cathode materials with high energy density for LIBs.11,12 In spite of these merits, however, there is no doubt that Lirich layered materials also suffer from several drawbacks, such as an initial irreversible capacity loss,13,14 poor cycling stability,15 and poor rate capability,16 which could impede their application in future electric vehicles. The reason for poor rate capability is partially due to the electronically insulated © 2018 American Chemical Society
Li2MnO3 phase denting the electron conductivity of Li-rich material.17,18 To improve the rate capability of these material, several different methods were used, such as minimizing particle size,19,20 carbon coating,21 and element doping.22 Among them, the substitution of metal elements such as V, Sn, Cr, Al, Fe, Mg, and Mo17,23−29 for transition elements in xLiMO2·(1−x)Li2MnO3 material has been proven to be effective at improving rate performance. Based on previous research, it has been shown that ionic and electronic conductivities of Li-rich material can be improved after cationic doping.24 Another effective improvement strategy is to enlarge the Li slab distance in the layered materials, thus improving the Li+ diffusion coefficient during the interaction/deinteraction process.30 Inspired by this idea, researchers introduced a large ionic radius of Na+ (0.102 nm) or K+ (0.138 nm) into the Li layer to enlarge the diffusion path and enhance the rate capability of Li-rich materials.31,32 At the same time, Li-rich materials with Li deficiency and layered/spinel heterostructure could significantly improve the rate capability, in which the spinel structure can offer 3D Li+ channels and then improve the Li+ diffusion coefficient.33−35 It is reported that severe Li deficiency in Li-rich materials may cause the Mn4+ in Li-rich materials to migrate, and layered-to-spinel transition can be Received: October 8, 2017 Revised: February 6, 2018 Published: February 18, 2018 4625
DOI: 10.1021/acssuschemeng.7b03634 ACS Sustainable Chem. Eng. 2018, 6, 4625−4632
Research Article
ACS Sustainable Chemistry & Engineering restrained during charge/discharge processes. Based on this method, Xin et al. reported that in the Li layer Ti4+ substitution of Li+ restricts the phase transition; the reason may be due to the fact that doped Ti4+ enhanced the activation energy of Mn migration, which hindered the Mn migration during the cycle.36 In this paper, we prepared Al-doped Li-rich materials Li1.5−xAlxMn0.675Ni0.1675Co0.1675O2 via a rapid nucleation and post-solvothermal method. After Al ions were introduced, the local environment of elements in the Li-rich material were changed and electrochemical performance was enhanced. To study the effect of Al3+ doping on the Li+ site in Li-rich materials, density functional theory (DFT) calculations are used.
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EXPERIMENTAL SECTION
Figure 1. XRD patterns of as-prepared samples.
Material Preparation. Li-rich layered material Li1.5−xAlxMn0.675Ni0.1675Co0.1675O2 (x = 0, 0.025, 0.05, 0.1, and 0.2, denoted as Al = 0, Al = 0.025, Al = 0.05, Al = 0.1, Al = 0.2, respectively) was synthesized by rapid nucleation and postsolvothermal method. A stoichiometric amount of NiSO4·6H2O, MnSO4·H2O, CoSO4·7H2O, and Al(NO3)3·9H2O was dissolved in a mixture solution (solution A), and then (NH4)2CO3 solution was poured into solution A rapidly. The suspension solution was transferred into a poly(tetrafluoroethylene) (Teflon) container and treated at 180 °C for 12 h to form precursors with spherical morphology. Then, these precursors were washed and dried in air at 60 °C for 12 h.37 To prepare the final cathode material, a desired amount of Li2CO3 was thoroughly mixed with the precursor. After being ground, the resulting mixture was calcined at 850 °C for 12 h in air. Characterizations. An X-ray diffractometer (Bruker AXS D8) equipped Cu Kα radiation was employed to obtain the X-ray diffraction (XRD) patterns. The morphology of as-prepared materials was analyzed by field emission scanning electron microscopy (Hitachi S-4800). A Thermo-DXR FT-Raman spectrometer with a Nd line laser source (λ = 532 nm) was used to collect Raman scattering of samples. X-ray photoelectron spectra (XPS) were obtained with PHI-5400 (PE, USA) equipment. Electrochemical Measurements. Electrochemical performance of as-prepared samples was evaluated by two-electrode cells in a 2032 coin cell. The composite cathode materials were prepared as follows: the synthesized samples, polyvinylidene fluoride (PVDF) binder, and conductive carbon (super-p) were mixed in a wt % ratio of 80:10:10 on Al foil, as described in previous literature.37 The cells were assembled in an Ar-filled glovebox, using 1 mol dm−3 LiPF6 in EC/ DMC (1:1 in volume) as electrolyte, lithium foil as the anode, and Celgard 2500 as separator. LANHE battery testers were used to measure the electrochemical performance in the voltage range between 2.0 and 4.8 V at 25 °C. Computational Details. The electronic structure calculations were performed with the spin-polarized DFT method using the Vienna Ab initio Simulation Package.38 Pseudopotentials used in the calculations were established by the projector-augmented wave method and the Perdew−Burke−Ernzerhof.39,40 Considering the strong correlative d electrons of transition metal Mn, the effective Hubbard-U parameter (Ueff = U − J = 4.9 eV) was employed.41 The plane-wave basis set was cut off by an energy of 500 eV. A 2 × 2 × 1 supercell (16 formula units) of Li2MnO3 with the space group symmetry C2/m was taken to build the configurations for the 6.25% doped Li2MnO3. The configurations were relaxed with a 2 × 2 × 2 centered k-mesh. A denser k-mesh (4 × 4 × 4) was used to obtain the density of states (DOS) of these relaxed configurations.
diffraction of NaFeO2 structure (R3̅m).13 These weak superlattice peaks were indexed to the monoclinic Li2 MnO3 component corresponding to the C2/m space group.42,43 The more Al3+ ions that were doped into bulk lattice of the layered structure, the more Li vacancies were formed because of the electrical neutralization; at the same time, the lattice of the material was expanded owing to the electrostatic repulsion, resulting in the (003) diffraction shifts from 18.695° in Al = 0 to 18.676° in Al = 0.2, as shown in Figure 1b. However, minor impurity peaks regarding the spinel phase can be observed in the patterns of Al = 0.1 and Al = 0.2 samples (Figure 1b), this is mainly due to the doping of Al3+ ions in the Li sites generating Li vacancies.36 Figure 2 shows the Rama spectrum of as-prepared samples. As can be seen, the Raman bands of pristine sample at 372 and
Figure 2. Raman spectrum of as-prepared samples.
422 cm−1 belong to the monoclinic Li2MnO3 component, and other obvious peaks at 477 and 588 cm−1 are associated with the bending Eg and stretching A1g modes, respectively.44−46 With the contents of Al increased from 0 to 0.2, Raman peaks of 372 and 422 cm−1 have disappeared, and 472 and 588 cm−1 were blue shifted toward a high wavenumber; this can be associated with an irreversible manganese migration into the interlayer lithium sites that results in the formation of a spineltype cation ordering.47,48 Meanwhile, we also found that their intensities decreased substantially; the reason may be that after Al doping the electronic conductivity of the material increases.45 For in-depth analysis of the relationship between doped ions and electrical conductivity of as-prepared materials, a potentiostatic polarization method was used, and the current dependence of polarization time was performed in Figure S1. The result exhibited that the Al = 0.2 sample has the highest electronic conductivity. In addition, the spectrum of Al-doped samples demonstrated a significant increase in the peak width of the Raman band at 599 and 611 cm−1; this could be an
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RESULTS AND DISCUSSION Figure 1 exhibits the XRD patterns of samples. Except for the weak superlattice reflection peaks arising between 20 and 25° (Figure 1a), most sample peaks can be assigned to the 4626
DOI: 10.1021/acssuschemeng.7b03634 ACS Sustainable Chem. Eng. 2018, 6, 4625−4632
Research Article
ACS Sustainable Chemistry & Engineering
ratio of Mn, Ni, Co, and Al was 0.675:0.166:0.163:0.057, approaching the theoretical value of the Al = 0.05 sample. To determine the oxidation states of each element and structures of as-prepared samples, the XPS spectra are shown in Figure 5. In the O 1s spectra, the peak located at 529.4 eV
indication of structural and lattice strain effects. Herein, we can draw a conclusion that the Al3+ ion doped into the bulk lattice of layered Li-rich material and changed the local coordination structure of the Li2MnO3 component. Figure 3 exhibits the SEM images of Al = 0 and Al = 0.05 samples. All samples exhibit a typical spherical morphology, and
Figure 5. XPS spectrum of (a) O 1s, (b) Al 2p, and (c) Mn 2p of asprepared samples.
corresponds to the lattice oxygen in the Li-rich material,23 and the peak around 531.4 eV can be assigned to the impurity Li oxides that exist on the surface of the compounds.49 As can be seen from Figure 5a, with the increase of Al doping, the gradually reduced peak intensity at 531.4 eV is assigned to the impurity oxides such as Li2O or Li2CO3, indicating that doped materials have more stability in air than in pristine material. The Al 2p peak located at 73.13 eV corresponds to that of LiAlO2 (73.4 eV) in the doped samples.50 As our previous work shows, the oxidation states of Ni, Co, Mn, and O are +2, +3, +4, and −2 in the Al = 0 sample.37 What we can see in Figure 5 and Figure S4 is that the peak positions of Mn, Co, Ni, and O all shift to higher energy positions with increased doping content of Al, which demonstrates that the chemical environment of these atoms has changed. At the same time, the additional peaks at 640.38 eV are observed in the spectrum of Al = 0.1 and Al = 0.2, indicating the formation of Mn3+. Due to the amount of Al3+ ion doped into the Li sites, some of transition metal may reduce the valence to maintain the electrical neutralization of as-prepared materials, thus minor Mn3+ may exist in the spectrum. Herein, we can draw a conclusion that Al3+ can dope the crystal structure and change the environment of the transition metal and oxygen atoms in the doped materials. To evaluate the influence of Al doping on the electrochemical performance of as-prepared samples, the first charge− discharge profiles of samples were conducted using a rate at 0.1 C, and the result is shown in Figure 6a. As normal Li-rich layered oxides, all electrodes show a sloping region and a long plateau; the former is attributed to Li+ deintercalation from the LiMn1/3Ni1/3Co1/3O2 component, and the latter corresponds to Li+ and O2− extracted from the Li2MnO3 phase and structural rearrangement in the initial charge process. First charge capacities of 117.4 mAh g−1, when voltage is less than 4.5 V, are observed in the Al = 0 sample. This value is nearly the same as the expected values based on the Ni2+/4+ and Co3+/4+ redox reaction (approximately 120 mAh g−1).51 With the increase of doped Al3+ content, charge profiles of as-prepared electrodes deliver higher capacities except for the Al = 0.2 electrode. This result indicates that the addition of Al has an influence on the crystal structure in the LiMn1/3Ni1/3Co1/3O2 component, and the reason for the higher charge capacities (
