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Hollow spherical lithium-rich layered oxide cathode material with suppressed voltage fading
Accepted Manuscript Hollow spherical lithium-rich layered oxide cathode material with suppressed voltage fading Weixiang Ding, Xueyang Cui, Jie Lei, Xiaodong Lin, Shengliang Zhao, Qi-Hui Wu, Mingsen Zheng, Quanfeng Dong PII:
S0013-4686(18)30124-5
DOI:
10.1016/j.electacta.2018.01.082
Reference:
EA 31064
To appear in:
Electrochimica Acta
Received Date: 14 November 2017 Revised Date:
28 December 2017
Accepted Date: 12 January 2018
Please cite this article as: W. Ding, X. Cui, J. Lei, X. Lin, S. Zhao, Q.-H. Wu, M. Zheng, Q. Dong, Hollow spherical lithium-rich layered oxide cathode material with suppressed voltage fading, Electrochimica Acta (2018), doi: 10.1016/j.electacta.2018.01.082. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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Graphic abstract (for review)
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Hollow spherical lithium-rich layered oxide cathode material with suppressed voltage fading
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Weixiang Ding,1,† Xueyang Cui,1,† Jie Lei,1 Xiaodong Lin,1 Shengliang Zhao,1 Qi-Hui Wu,*2 Mingsen Zheng*1 and Quanfeng Dong*1 1
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Collaborative Innovation Centre of Chemistry for Energy Materials (iChEM), State Key Laboratory of
Physical Chemistry of Solid Surfaces, Department of Chemistry, College of Chemistry and Chemical Engineering, Xiamen University, Xiamen, Fujian 361005, China. 2
Department of Materials Chemistry, School of Chemical Engineering and Materials Science, Quanzhou
Normal University, Quanzhou, Fujian 362000, China. These authors contributed equally to this work.
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†
*Corresponding author e-mail: [email protected]; [email protected]; [email protected]
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Abstract
Lithium-rich layered oxide materials are very promising but with some weaknesses, which
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limit their commercialization, such as large irreversible capacity loss at the first cycle, continuing discharge capacity and voltage fading. Among these, voltage fading is one of the most serious problems for their final practical applications. Herein, we successfully synthesized hollow spherical Li1.2Ni0.13Co0.13Mn0.54O2 cathode material with MnO2 hollow spheres as the templates. The as-prepared sample was spherically secondary particles with hollow cavities. It showed improved cycle stability, rate capability, and suppression of voltage fading. After 100 cycles at 1 C, the retention of average discharge voltage could reach up to 97.0 % (the average voltage loss
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was only 0.103 V). It can be attributed to the hollow structure which could release the stress of lithium removal, leading to the suppressed migration of transition metal ions into the lithium layers and the transformation from layered phase to spinel-like phase.
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Keywords：lithium-ion batteries, Li-rich layered oxides, voltage fading, phase transition, micro hollow structure
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1. Introduction
Since commercialized by Sony Corporation, lithium ion battery (LIB), as the power supply
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for electronic devices, has been an integral part of our lives, and drew much attentions across the world due to its high energy density and long service life [1]. With the development of electrical vehicles (EVs) and hybrid electrical vehicles (HEVs) and smart grids, the promotion of LIB cathode materials is the main key to meet the high requirements for these arising applications [2-
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4]. Currently, the 3C market (computer, communication and consumer electronic) has been filled with LiCoO2, and the layered LiMO2 (M=Ni、Co and Mn), spinel LiMn2O4 and olivine-type LiFePO4, but they are not to satisfy the practical requirements in the fields of EVs and HEVs due
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to their low capacity and energy density [5-7]. The lithium-rich layered transition metal oxides
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(Li2MnO3-LiMO2, M=Ni, Co, Mn) possess large reversible capacity (> 250 mAh·g-1) and relatively low cost， which are regarded as promising candidates for high energy density LIBs [8-10]. However, they have some considerable defects, such as the large irreversible capacity loss at the first cycle [11-13], continuous discharge capacity and voltage fading during cycling, resulted from the structural reorganization of materials affected by both the erosion of electrolyte and the phase transformation during the charge-discharge process [14-17]. Among all those drawbacks mentioned above, the problem of voltage fading has particularly
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bedeviled the researchers. There have already been numerous previous reports focused on the mechanism of voltage fading. According to Wang et al.’s work, the irreversible phase transition is the main and root reason of voltage fading [16]. Daniel et al. have confirmed that the ions of
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transition metals transferred their position from the octahedral transition metal layer to the octahedral lithium (Li) layer, and the irreversible position transfer of Li from octahedron to tetrahedron are the major mechanism of the irreversible phase transition [18]. In order to solve
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the problem of voltage fading for Li-rich materials, researchers have dedicated massive efforts. Zhang et al. have adopted the guar gum as the binder of Li-rich cathode materials for the first
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time. The decrease of average voltage for the cells with PVDF and guar gum binders is 0.61 V and 0.35 V after 200 cycles at 100 mA·g-1 between 2.0 and 4.8 V, respectively [19]. Ma, et al. decelerated voltage decay by selenium doping and the cathode material exhibited a mid-point voltage (MPV) retention of 95% after 100 cycles at 0.1 C between 2.5 and 4.6 V [20]. Chong, et
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al. suppressed voltage decay of Li-rich cathode material by a surface coating of CoF2. 1 % CoF2coated sample exhibits an ameliorative voltage drop of 0.312 V at 0.1 C after 100 cycles between 2.0 and 4.8 V [21].
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In this paper, the micrometer hollow spherical Li1.2Ni0.13Co0.13Mn0.54O2 cathode material was synthesized by using a simple and feasible sacrificing template method. The structure,
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morphology and electrochemical performance were characterized by XRD, SEM, XPS and electrochemical technique. We paid more attention to the voltage fade of the hollow cathode material after long cycles. It was interesting to find that the hollow cathode delivered a high average voltage retention (97.0 % after 100 cycles and 93.5 % after 200 cycles) at 1 C. Compared to that of the infarctate spherical Li-rich material, the voltage decay is rather small. Then we used XRD and XPS technology to verify that the hollow spherical structure could
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suppress the irreversible transformation from layered phase to spinel-like phase, which may be the main reason retarding the fading of cell potential.
2.1 Synthesis of hollow spherical Li1.2Ni0.13Co0.13Mn0.54O2
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2. Experimental
5.7 mmol (0.963 g) of MnSO4 and 5.7 mmol (0.451 g) of NH4HCO3 were dissolved in 100 ml
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of a water/ethanol mixed solution (volume ratio of 9:1), respectively. Then the NH4HCO3 solution was added dropwise into the MnSO4 solution under vigorous stirring at room
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temperature. After completion of the reaction, the mixture was filtered, washed with water and ethanol for several times, and dried in an oven at 80 oC to obtain the pale pink MnCO3 powders. The powders were then treated at 400 oC for 5 h in air at a temperature increasing rate of 1 oC min-1. After cooling down the furnace naturally the black MnO2 powders were gained. The obtained hollow spherical MnO2 precursor was then mixed with Ni(NO3)2·6H2O,
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Co(NO3)2·6H2O, MnO2, LiOH·H2O, and NaCl at the molar ratio of 0.24: 0.24: 1: 2.33: 7.0, then was dispersed into 50 mL of acetone, stirring slowly at 50 oC under the conditions of evaporation. The powders were scraped off and ground, then calcined in air at 850 oC for 15 h. The solid thus
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obtained was washed with water for three times to remove the NaCl and then washed with
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ethanol and dried in an oven at 80 oC to obtain hollow sphere Li1.2Ni0.13Co0.13Mn0.54O2 (HLRMS). The infarctate sphere Li1.2Ni0.13Co0.13Mn0.54O2 (ILRMS) was synthesized as well used as a contrast sample. MnCO3 powders were treated at 500 oC in air for 20 min, with a heating rate of 6 oC min-1. After obtaining the infarctate sphere MnO2 powders, they were mixed with metal salt and calcined using the same method as HLRMS to finally get ILRMS. 2.2 Materials characterizations The crystal structure of the samples was characterized using X-ray powder diffraction (XRD,
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Rigaku Ultima IV X-ray diffractormeter) with Cu Kα1 radiation (λ=1.5405 Å). The diffraction data were obtained at 2θ=10°-80°, with a step size of 0.02°. The morphology of samples was obtained from field-emission scanning electron microscopy (SEM, Hitachi S-4800, 15 kV)
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equipped with an energy dispersive X-ray analysis (EDX) system. X-ray photoelectron spectroscopy (XPS) was collected at PHI 5000 VP III instruments using Al Kα source. 2.3 Electrochemical measurements
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Electrochemical performances were evaluated with CR2016 coin cells between 2.0-4.8 V. We made the cathodes for the battery test cells from the active material, Super P and water solute
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polymer n-lauryl acrylate (LA Chengdu, China) in the weight ration of 8:1:1. Al foil was used as the current collector and Li foil was used as the counter electrode. The electrolyte consisted of a solution of 1.0 M LiPF6 in propylene carbonate (PC)/ethyl methyl carbonate (EMC)/sulfolane (SL) (1:1:1, in volume). And the separator was a Celgard 2400 polypropylene membrane.
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Cell assembly was carried out in an Ar-filled glove box with the concentration of moisture and oxygen below 1 ppm. The cells were galvanostatically charged and discharged using a BTS Battery Tester (Neware, Shenzhen, China).
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3. Results and discussion
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The morphology and particle size of the MnO2 precursors were examined using SEM. Fig. 1c and 1d show the images of infarctate MnO2 precursor. The sample has a spherical morphology with secondary particle size about 2-3 µm. It can be seen from few broken particles that the center of the secondary particles is solid. Similar to the solid MnO2 precursor, the secondary particle of the hollow MnO2 is also sphere with a size of about 2-4 µm (Fig. 1a and b). We can directly identify the hollow structure by a few broken secondary particles. According to the Kirkendall effect [22, 23], during the calcining process, the metal atoms rapidly diffuse outward
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and the opposite O atoms slowly diffuse inward, which leads to the formation of the hole in MnO2 precursor. When the calcination temperature increases, the diffusion rate of all atoms tends to be consistent leading to the disappearance of the hole.
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In Fig. 2, the diffraction peaks of two electrode materials could be attributed to α-NaFeO2 structure (space group R-3m), which do not show any obvious differences. Weak diffractions could be seen at 2θ = 20 ~ 25°, which is Li2MnO3 in the transition metal layer to form Li and Mn
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ordered arrangement of superlattice peak [24-26]. It could be seen that the diffraction peaks of (006)/(012) and (018)/(110) had obvious cleavages, indicating that each synthesized material has
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a good layered structure [27]. It has been reported that the ordering of the structure can be indicated from the I(003)/I(004) and the c/a [28]. As seen from the Table 1, the I(003)/I(104) and the c/a of HLRMS are both greater than those of ILRMS, which means that HLRMS has a better ordering and ion arrangement. We owe it to the cavity of MnO2 precursor. The cavity increases
calcination.
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the contact surface when mixed with metal salt and shorten the ion diffusion path when
Both the as-prepared HLRMS (Fig. 3a, b) and ILRMS (Fig. 3c, d) were spherical particles.
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The hollow cavity existed in the cross section of HLRMS spherical particle with a diameter about 2 µm. The secondary particle radius was about 3 µm and the wall thickness was about 2
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µm. ILRMS, with radius of about 3 µm, have no obvious cavity which could be seen. Hollow structure of HLRMS was conducive to the infiltration of the electrolyte and rapid reaction, and buffered the volume change of material in the charge and discharge process. Especially, the particles cracking and erosion caused by the difference of the volume change between the different components (Li2MnO3 and LiMO2). The elemental ratios obtained by EDX analysis are shown in Fig. 4a. It could be seen that the
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materials contained Ni, Co and Mn transition metal elements, and their contents were similar to those of the formulas we designed. Fig. 4b, c and d are XPS spectra of these three elements. The Mn 2p spectrum has two main peaks at about 642.6 and 654.1 eV, corresponding to Mn 2p3/2 and
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Mn 2p1/2, respectively. The spin-orbital splitting energy is about 11.5 eV, the value matches well with the binding energy of Mn4+ [29, 30]. Ni 2p3/2 is located at 854.5 eV and Ni 2p1/2 is at 872.4 eV. Ni 2p3/2 and 2p1/2 with spin-level splitting of 17.9eV indicates that the element is in the form
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of Ni2+ [31-33]. The two emission lines of Co 2p are at about 780.0 eV and 795.5 eV, indicating the existence of Co3+ [32-35]. Therefore, the XPS spectra suggest the oxidation states of Ni, Co
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and Mn to be 2+, 3+ and 4+, respectively.
As a kind of LIB cathode materials, Li-rich manganese base materials have a huge advantage in the specific capacity. As we can see from Fig. 5a, the charge and discharge curves of both two materials show typical characteristics of the layered Li-rich material. The plateau at 4.5 V in the
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first cycle was attributed to the activation of the Li2MnO3, which lose oxygen and form electrochemical activity layered MnO2. HLRMS delivers charge and discharge capacity of 321.1 and 268.7 mAh·g-1 respectively, while ILRMS shows charge and discharge capacity of 323.9 and
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262.6 mAh·g-1. The Coulombic efficiencies of two materials are 83.7 % and 81.7 %. The HLRMS charge platform is lower than that of ILRMS, and the discharge platform is higher.
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Thus, we can conclude that the hollow structure of the material can make the reduction of polarization during charge and discharge. Li+ ion diffusions were thus easy due to the hollowly secondary structure, which is advantageous to the infiltration of the electrolyte and rapid reaction.
The cycling performance of HLRMS and ILRMS cathodes tested at a rate of 1 C in the potential range of 2.0-4.8 V is shown in Fig. 5b. It can be seen that the discharge capacity of
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HLRMS has almost not declined after 100 cycles at 1 C. After 200 cycles, discharge capacities of 174.5 mAh·g-1 and 111.6 mAh·g-1 were retain respectively for HLRMS and ILRMS, which is 94.7 % and 58.9 % of the first discharge capacity. These results reflect the better stability of
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HLRMS cathode over cycling.
The battery based on HLRMS was cycled at different rates ranging from 0.2 C to 5 C (Fig. 5c). The HLRMS cathode delivers a discharge capacity of 243.7, 206.6, 184.2, 165.7, 138.1
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mAh·g-1 at rate of 0.2 C, 0.5 C, 1 C, 2 C and 5 C respectively. We also compared the rate performance of the HLRMS with that of ILRMS at different rates. The battery based on ILRMS
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was examined under the same conditions. The ILRMS delivers a discharge capacity of 227.1, 184.6, 162.0, 134.3, 93.8 mAh·g-1 at rate of 0.2 C, 0.5 C, 1 C, 2 C and 5 C. We can conclude that the HLRMS shows an obvious improvement on the rate performance. In order to illustrate the advantages of the hollow micro-sphere particles, the morphologies of
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HLRMS and ILRMS after cycling are presented (Fig. 6). These two samples before cycling are both sphere particles with diameter about 6 µm (Fig. 3). After cycling, the spherical particles of HLRMS still can be identified on the electrode (Fig. 6a, b), which indicates the morphology of
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the material barely changed and the structure of HLRMS is stable, while the original morphology of ILRMS can’t be found as indicated in Fig. 6c and d. The hollow structure releases the stress of
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Li+ intercalation/ deintercalation and buffers the volume change of material, thus avoiding the cracking of particles. The destruction of particles leads to increase of contact interface between electrode material and electrolyte. The side reactions of active material and electrolyte result in capacity fading and voltage decay.[36] We can see from the Fig. S2b that the material turned into spinel phase and rock-salt phase in the outer surface, only small fraction of layered phase remained. Spacings of the lattice are ~0.47 nm which approximate to the (003) interplanar
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spacing of the Li-rich layered material. And the transformation from layered phase to spinel phase can also be seen from the exposed original inner region (Fig. S2c and d). This illustrates that the original inner region is exposed and transform into spinel-like phase, which causes the
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deepening of voltage decay.
The visual representation of the discharge voltage drop during the cycle is given in Fig. 7b and 7c. After the discharge capacity normalized, the discharge curves of the materials change in
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different cycles could be seen more clearly. It could be observed from the comparison that the voltage fading of HLRMS is much less than that of ILRMS. The average voltage of the HLRMS
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is 3.49 V at the rate of 1 C. The average voltage retention of HLRMS is 97.0 % after 100 cycles and 93.5 % after 200 cycles (Fig. 7a). The average voltage of ILRMS is 3.48 V. The average voltage retention of ILRMS is 91.3 % after 100 cycles and 86.3 % after 200 cycles. Fig. 7(d, e, f) shows the XRD patterns of the material before and after 100 cycles at 1 C. If the material has
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spinel-like phase, we could observe the peak of the hetero phase before the (104) diffraction peak of the layered phase. It can be seen from the (104) diffraction peak in the XRD pattern that a new diffraction peak appears at about 2θ = 44.2° after charge and discharge of the material, which
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belongs to the spinel phase (440) plane, indicating that transformation of layered phase to spinellike phase occurs after cycling [37]. The spinel phase peak (Isp(440)) of the ILRMS is found to be
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stronger than that of the HRLMS, it means that the spinel-like phase transition of the ILRMS material is more severe than that of the HLRMS after charging/discharging. The charge-discharge process of the Li-rich material generally includes the reactions shown below (E1-E6). The remove of Li+ ions in the layered phase during the charge process is accompanied by the oxidation of the transition metal ions (E1).[38, 39] When the voltage reached about 4.5 V, the activation of Li2MnO3 takes place (E2), and Li+ and O2 are removed to
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form electrochemical active layered MnO2. The process is irreversible and is closely related to the characteristics of high capacity and voltage fading.[40] As shown in the differential capacity curve of the material in Fig. 8, reduction of Co ions occurrs at about 4.2 V (E3), reduction of Ni
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ions happens at about 3.7 V (E4),[41] the inserting of Li+ ions into MnO2 takes place at about 3.3 V (E5) and Li intercalation into the spinel-like phase is at about 3.0 V (E6).[42] As shown in Fig. 8b, the capacity at about 3.0 V increases continuously and the capacity above 3.5 V greatly
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reduces, which implies that the capacity could be assigned to the layered phase decreases and some layered phase transforming into spinel-like phase. The capacity about 3.0 V of HLRMS
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(Fig. 8a) barely increases after 20 cycles and the capacity above 3.5 V is more stable than that of ILRMS (Fig. 8b). Charge: LiMO2→Li++MO2
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Li2MnO3→Li++MnO2+O2 Discharge: Co4+→Co3+ Ni4+→Ni2+
E1 E2
E3 E4 E5
MnO2+xLi+→LixMnO2→LiMn2O4 (spinel)→Li2Mn34+O7(spinel)+Mn2+O
E6
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MnO2+Li+→LiMn3+O2
The spinel phase transformation of Lithium-rich layered cathode materials contains the following specific processes. In charging process, metal ions (Co3+ and Ni2+) are oxidized to higher oxidation valences. With some Ni ions entering the Li layer, the Li-rich spinel Li(Li1/3Ni1/2Mn7/6)O4 structure was mainly formed on the surface region [14]. In discharging process, Ni ions at the TM layer are reduced from higher oxidation valences to +2 valence, but it
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is difficult to restore the Ni ions to the original state [30]. When Li ions are intercalated into MnO2, Mn ions are reduced to +3 valence, while the Mn ions of layered phase remain unchanged (+4 valence). Then the LixMnO2 tended to transform into spinel-like phase LiMn2O4, which in
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turn formed Li2Mn34+O7 and MnO (E6) further by disproportionation. Consequently, the MnO turns into fragments and stay in the electrolyte, resulting in capacity loss of the material (~3.3 V) [43]. Because of these reactions, the spinel phase transformation can be indirectly detected with
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XPS.
Fig. 9 shows XPS diagrams of two materials (HLRMS and ILRMS) before and after charge
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and discharge. In Li1.2Ni0.13Co0.13Mn0.54O2, Ni ions are of +2 valence (~855 eV), and Mn ions are of +4 valence (~642.6 eV). After charge and discharge of HLRMS, Ni 2p3/2 major peak will slightly move to +3 valence (~857.6 eV), and Mn elements of +3 valence (~641.1 eV) and of +4 valence (~642.6 eV) coexist, showing that the valence state of the elements agrees with that
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mentioned above. A new Ni 2p3/2 peak centered at ~849.6 eV is caused by the Ni ions at the Li layer [44]. After charge and discharge of ILRMS, the fraction of Ni3+ ions is much higher than that of HLRMS, and the peak of 849.6 eV is stronger, indicating that more Ni ions of ILRMS
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enter the Li layer on the surface to generate spinel-like phase. In addition, Mn elements of both +3 valence and of +4 valence exist. However, comparing with HLRMS, the content of Mn4+ ions
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is relatively high, and that of Mn3+ elements is rather lower, which agrees with the spinel phase transformation and disproportionated reaction of E6, which indicates the spinel phase transformation of ILRMS is more intense. Meanwhile, it can be seen that the displacement of Ni 2p3/2 major peak of ILRMS towards high valence is significant, which tallies with the fact that the capacity fade of the material is significant. Besides, the Ni elements of high valence on material surface are liable to have side reaction with electrolyte to further deteriorate the
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electrochemical performance of the material. The suppression of phase transformation for HLRMS is attributed to two points: (1) stable secondary particles. The secondary particle may collapse after long term cycles, which lead to the increase of contact surface between electrode
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and electrolyte. The side reaction of active material and electrolyte results in the electrode polarization associate with the voltage fading during charge and discharge. HLRMS can keep its morphology after charge and discharge, which own to cavity with appropriate size. (2) The
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phenomenon of the transformation from layered phase to spinel phase involve microstress during Li intercalation/deintercalation. Stress of Li migration will engendered the migration of transition
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metal ions into the Li layers and the formation of spinel phase [37, 45, 46]. The stress caused by Li+ intercalation/deintercalation can be released to avoid the structure destruction and the spinel phase formation. 4. Conclusions
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Hollow spherical layered Li-rich cathode material Li1.2Ni0.13Co0.13Mn0.54O2 was successfully synthesized via a facile lithiation method by using MnO2 hollow spheres as template. The cathode material HLRMS exhibits a favorable enhancement of cycle stability compared with the
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control group ILRMS, especially in the suppression of voltage fading. The capacity of the
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cathode material reaches up to 94.7 % after 200 cycles at 1 C and the average discharge voltage remains 97.0% after 100 cycles and 93.5% after 200 cycles at 1 C. The improvements of the electrochemical performance can be attributed to the hollow secondary particle structure, which release the stress caused by Li+ intercalation/deintercalation. Thus, the structure stable is guaranteed and transformation from layered phase to spinel phase is suppressed. The demonstration of the Li-rich cathode material with hollow spherical particles could offer an efficient way to suppress the voltage decay and to promote its practical applications.
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Acknowledgments This work was supported by the National 973 Program (2015CB251102), the Key Project of NSFC (U1305246, 21673196, 21621091), and the Fundamental Research Funds for the Central
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Table 1 The lattice and structural parameters of HLRMS and ILRMS. c(Å)
c/a
I(003)/I(104)
HLRMS
2.8478
14.21321
4.9909
2.320
ILRMS
2.84944
14.19371
4.9812
2.214
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a(Å)
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All Figure captions
Fig. 1 SEM images of (a, b) the hollow MnO2 precursor and (c, d) the infarctate
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MnO2 precursor.
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Fig. 2 XRD patterns of HLRMS and ILRMS.
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Fig. 3 SEM images of (a, b) HLRMS, (c, d) ILRMS.
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Fig. 4 (a) Elemental ratios measured by EDX and (b, c, d) XPS spectra of HLRMS and ILRMS.
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Fig. 5 Electrochemical performances of HLRMS and ILRMS. (a) the charge and discharge curves at 0.1 C, (b) the cycle performance tested at 1 C, (c) rate performance (1 C=200 mA·g-1).
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Fig. 6 SEM images of (a, b) HLRMS and (c, d) ILRMS after 100 cycles (at 1C).
Fig. 7 (a) Average voltage of HLRMS and ILRMS tested at 1C, (b, c) normalized
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discharge capacity curves of HLMRS and ILRMS, (d, e) XRD patterns of the electrodes (before and after 100 cycles at 1C) and (f) the enlarge XRD pattern between 43-46° (the asterisked peaks represent the diffraction peaks of Al foil).
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Fig. 8 Differential curves (dQ/dV) of (a) HLRMS and (b) ILRMS.
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Fig. 9 XPS patterns of the electrodes (before and after 100 cycles at 1 C).
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Regulation on the secondary particle hollow structure enables Li1.2Ni0.13Co0.13Mn0.54O2 to serve as cathode electrode for high power lithium ion battery showing excellent electrochemical performance.
It proves that the hollow structure of material is really favorable for material to keep the
It demonstrates that the morphology and structure of rich lithium material exert significant impacts on the voltage attenuation of material. The hollow structure can obviously inhibit the voltage, because it can effectively suppress the transformation from layered phase to
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spinel phase.
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original morphology and structure, greatly improving the cycling stability of the material.
S0013-4686(18)30124-5
DOI:
10.1016/j.electacta.2018.01.082
Reference:
EA 31064
To appear in:
Electrochimica Acta
Received Date: 14 November 2017 Revised Date:
28 December 2017
Accepted Date: 12 January 2018
Please cite this article as: W. Ding, X. Cui, J. Lei, X. Lin, S. Zhao, Q.-H. Wu, M. Zheng, Q. Dong, Hollow spherical lithium-rich layered oxide cathode material with suppressed voltage fading, Electrochimica Acta (2018), doi: 10.1016/j.electacta.2018.01.082. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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Graphic abstract (for review)
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Hollow spherical lithium-rich layered oxide cathode material with suppressed voltage fading
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Weixiang Ding,1,† Xueyang Cui,1,† Jie Lei,1 Xiaodong Lin,1 Shengliang Zhao,1 Qi-Hui Wu,*2 Mingsen Zheng*1 and Quanfeng Dong*1 1
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Collaborative Innovation Centre of Chemistry for Energy Materials (iChEM), State Key Laboratory of
Physical Chemistry of Solid Surfaces, Department of Chemistry, College of Chemistry and Chemical Engineering, Xiamen University, Xiamen, Fujian 361005, China. 2
Department of Materials Chemistry, School of Chemical Engineering and Materials Science, Quanzhou
Normal University, Quanzhou, Fujian 362000, China. These authors contributed equally to this work.
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†
*Corresponding author e-mail: [email protected]; [email protected]; [email protected]
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Abstract
Lithium-rich layered oxide materials are very promising but with some weaknesses, which
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limit their commercialization, such as large irreversible capacity loss at the first cycle, continuing discharge capacity and voltage fading. Among these, voltage fading is one of the most serious problems for their final practical applications. Herein, we successfully synthesized hollow spherical Li1.2Ni0.13Co0.13Mn0.54O2 cathode material with MnO2 hollow spheres as the templates. The as-prepared sample was spherically secondary particles with hollow cavities. It showed improved cycle stability, rate capability, and suppression of voltage fading. After 100 cycles at 1 C, the retention of average discharge voltage could reach up to 97.0 % (the average voltage loss
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was only 0.103 V). It can be attributed to the hollow structure which could release the stress of lithium removal, leading to the suppressed migration of transition metal ions into the lithium layers and the transformation from layered phase to spinel-like phase.
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Keywords：lithium-ion batteries, Li-rich layered oxides, voltage fading, phase transition, micro hollow structure
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1. Introduction
Since commercialized by Sony Corporation, lithium ion battery (LIB), as the power supply
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for electronic devices, has been an integral part of our lives, and drew much attentions across the world due to its high energy density and long service life [1]. With the development of electrical vehicles (EVs) and hybrid electrical vehicles (HEVs) and smart grids, the promotion of LIB cathode materials is the main key to meet the high requirements for these arising applications [2-
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4]. Currently, the 3C market (computer, communication and consumer electronic) has been filled with LiCoO2, and the layered LiMO2 (M=Ni、Co and Mn), spinel LiMn2O4 and olivine-type LiFePO4, but they are not to satisfy the practical requirements in the fields of EVs and HEVs due
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to their low capacity and energy density [5-7]. The lithium-rich layered transition metal oxides
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(Li2MnO3-LiMO2, M=Ni, Co, Mn) possess large reversible capacity (> 250 mAh·g-1) and relatively low cost， which are regarded as promising candidates for high energy density LIBs [8-10]. However, they have some considerable defects, such as the large irreversible capacity loss at the first cycle [11-13], continuous discharge capacity and voltage fading during cycling, resulted from the structural reorganization of materials affected by both the erosion of electrolyte and the phase transformation during the charge-discharge process [14-17]. Among all those drawbacks mentioned above, the problem of voltage fading has particularly
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bedeviled the researchers. There have already been numerous previous reports focused on the mechanism of voltage fading. According to Wang et al.’s work, the irreversible phase transition is the main and root reason of voltage fading [16]. Daniel et al. have confirmed that the ions of
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transition metals transferred their position from the octahedral transition metal layer to the octahedral lithium (Li) layer, and the irreversible position transfer of Li from octahedron to tetrahedron are the major mechanism of the irreversible phase transition [18]. In order to solve
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the problem of voltage fading for Li-rich materials, researchers have dedicated massive efforts. Zhang et al. have adopted the guar gum as the binder of Li-rich cathode materials for the first
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time. The decrease of average voltage for the cells with PVDF and guar gum binders is 0.61 V and 0.35 V after 200 cycles at 100 mA·g-1 between 2.0 and 4.8 V, respectively [19]. Ma, et al. decelerated voltage decay by selenium doping and the cathode material exhibited a mid-point voltage (MPV) retention of 95% after 100 cycles at 0.1 C between 2.5 and 4.6 V [20]. Chong, et
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al. suppressed voltage decay of Li-rich cathode material by a surface coating of CoF2. 1 % CoF2coated sample exhibits an ameliorative voltage drop of 0.312 V at 0.1 C after 100 cycles between 2.0 and 4.8 V [21].
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In this paper, the micrometer hollow spherical Li1.2Ni0.13Co0.13Mn0.54O2 cathode material was synthesized by using a simple and feasible sacrificing template method. The structure,
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morphology and electrochemical performance were characterized by XRD, SEM, XPS and electrochemical technique. We paid more attention to the voltage fade of the hollow cathode material after long cycles. It was interesting to find that the hollow cathode delivered a high average voltage retention (97.0 % after 100 cycles and 93.5 % after 200 cycles) at 1 C. Compared to that of the infarctate spherical Li-rich material, the voltage decay is rather small. Then we used XRD and XPS technology to verify that the hollow spherical structure could
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suppress the irreversible transformation from layered phase to spinel-like phase, which may be the main reason retarding the fading of cell potential.
2.1 Synthesis of hollow spherical Li1.2Ni0.13Co0.13Mn0.54O2
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2. Experimental
5.7 mmol (0.963 g) of MnSO4 and 5.7 mmol (0.451 g) of NH4HCO3 were dissolved in 100 ml
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of a water/ethanol mixed solution (volume ratio of 9:1), respectively. Then the NH4HCO3 solution was added dropwise into the MnSO4 solution under vigorous stirring at room
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temperature. After completion of the reaction, the mixture was filtered, washed with water and ethanol for several times, and dried in an oven at 80 oC to obtain the pale pink MnCO3 powders. The powders were then treated at 400 oC for 5 h in air at a temperature increasing rate of 1 oC min-1. After cooling down the furnace naturally the black MnO2 powders were gained. The obtained hollow spherical MnO2 precursor was then mixed with Ni(NO3)2·6H2O,
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Co(NO3)2·6H2O, MnO2, LiOH·H2O, and NaCl at the molar ratio of 0.24: 0.24: 1: 2.33: 7.0, then was dispersed into 50 mL of acetone, stirring slowly at 50 oC under the conditions of evaporation. The powders were scraped off and ground, then calcined in air at 850 oC for 15 h. The solid thus
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obtained was washed with water for three times to remove the NaCl and then washed with
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ethanol and dried in an oven at 80 oC to obtain hollow sphere Li1.2Ni0.13Co0.13Mn0.54O2 (HLRMS). The infarctate sphere Li1.2Ni0.13Co0.13Mn0.54O2 (ILRMS) was synthesized as well used as a contrast sample. MnCO3 powders were treated at 500 oC in air for 20 min, with a heating rate of 6 oC min-1. After obtaining the infarctate sphere MnO2 powders, they were mixed with metal salt and calcined using the same method as HLRMS to finally get ILRMS. 2.2 Materials characterizations The crystal structure of the samples was characterized using X-ray powder diffraction (XRD,
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Rigaku Ultima IV X-ray diffractormeter) with Cu Kα1 radiation (λ=1.5405 Å). The diffraction data were obtained at 2θ=10°-80°, with a step size of 0.02°. The morphology of samples was obtained from field-emission scanning electron microscopy (SEM, Hitachi S-4800, 15 kV)
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equipped with an energy dispersive X-ray analysis (EDX) system. X-ray photoelectron spectroscopy (XPS) was collected at PHI 5000 VP III instruments using Al Kα source. 2.3 Electrochemical measurements
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Electrochemical performances were evaluated with CR2016 coin cells between 2.0-4.8 V. We made the cathodes for the battery test cells from the active material, Super P and water solute
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polymer n-lauryl acrylate (LA Chengdu, China) in the weight ration of 8:1:1. Al foil was used as the current collector and Li foil was used as the counter electrode. The electrolyte consisted of a solution of 1.0 M LiPF6 in propylene carbonate (PC)/ethyl methyl carbonate (EMC)/sulfolane (SL) (1:1:1, in volume). And the separator was a Celgard 2400 polypropylene membrane.
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Cell assembly was carried out in an Ar-filled glove box with the concentration of moisture and oxygen below 1 ppm. The cells were galvanostatically charged and discharged using a BTS Battery Tester (Neware, Shenzhen, China).
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3. Results and discussion
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The morphology and particle size of the MnO2 precursors were examined using SEM. Fig. 1c and 1d show the images of infarctate MnO2 precursor. The sample has a spherical morphology with secondary particle size about 2-3 µm. It can be seen from few broken particles that the center of the secondary particles is solid. Similar to the solid MnO2 precursor, the secondary particle of the hollow MnO2 is also sphere with a size of about 2-4 µm (Fig. 1a and b). We can directly identify the hollow structure by a few broken secondary particles. According to the Kirkendall effect [22, 23], during the calcining process, the metal atoms rapidly diffuse outward
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and the opposite O atoms slowly diffuse inward, which leads to the formation of the hole in MnO2 precursor. When the calcination temperature increases, the diffusion rate of all atoms tends to be consistent leading to the disappearance of the hole.
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In Fig. 2, the diffraction peaks of two electrode materials could be attributed to α-NaFeO2 structure (space group R-3m), which do not show any obvious differences. Weak diffractions could be seen at 2θ = 20 ~ 25°, which is Li2MnO3 in the transition metal layer to form Li and Mn
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ordered arrangement of superlattice peak [24-26]. It could be seen that the diffraction peaks of (006)/(012) and (018)/(110) had obvious cleavages, indicating that each synthesized material has
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a good layered structure [27]. It has been reported that the ordering of the structure can be indicated from the I(003)/I(004) and the c/a [28]. As seen from the Table 1, the I(003)/I(104) and the c/a of HLRMS are both greater than those of ILRMS, which means that HLRMS has a better ordering and ion arrangement. We owe it to the cavity of MnO2 precursor. The cavity increases
calcination.
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the contact surface when mixed with metal salt and shorten the ion diffusion path when
Both the as-prepared HLRMS (Fig. 3a, b) and ILRMS (Fig. 3c, d) were spherical particles.
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The hollow cavity existed in the cross section of HLRMS spherical particle with a diameter about 2 µm. The secondary particle radius was about 3 µm and the wall thickness was about 2
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µm. ILRMS, with radius of about 3 µm, have no obvious cavity which could be seen. Hollow structure of HLRMS was conducive to the infiltration of the electrolyte and rapid reaction, and buffered the volume change of material in the charge and discharge process. Especially, the particles cracking and erosion caused by the difference of the volume change between the different components (Li2MnO3 and LiMO2). The elemental ratios obtained by EDX analysis are shown in Fig. 4a. It could be seen that the
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materials contained Ni, Co and Mn transition metal elements, and their contents were similar to those of the formulas we designed. Fig. 4b, c and d are XPS spectra of these three elements. The Mn 2p spectrum has two main peaks at about 642.6 and 654.1 eV, corresponding to Mn 2p3/2 and
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Mn 2p1/2, respectively. The spin-orbital splitting energy is about 11.5 eV, the value matches well with the binding energy of Mn4+ [29, 30]. Ni 2p3/2 is located at 854.5 eV and Ni 2p1/2 is at 872.4 eV. Ni 2p3/2 and 2p1/2 with spin-level splitting of 17.9eV indicates that the element is in the form
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of Ni2+ [31-33]. The two emission lines of Co 2p are at about 780.0 eV and 795.5 eV, indicating the existence of Co3+ [32-35]. Therefore, the XPS spectra suggest the oxidation states of Ni, Co
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and Mn to be 2+, 3+ and 4+, respectively.
As a kind of LIB cathode materials, Li-rich manganese base materials have a huge advantage in the specific capacity. As we can see from Fig. 5a, the charge and discharge curves of both two materials show typical characteristics of the layered Li-rich material. The plateau at 4.5 V in the
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first cycle was attributed to the activation of the Li2MnO3, which lose oxygen and form electrochemical activity layered MnO2. HLRMS delivers charge and discharge capacity of 321.1 and 268.7 mAh·g-1 respectively, while ILRMS shows charge and discharge capacity of 323.9 and
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262.6 mAh·g-1. The Coulombic efficiencies of two materials are 83.7 % and 81.7 %. The HLRMS charge platform is lower than that of ILRMS, and the discharge platform is higher.
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Thus, we can conclude that the hollow structure of the material can make the reduction of polarization during charge and discharge. Li+ ion diffusions were thus easy due to the hollowly secondary structure, which is advantageous to the infiltration of the electrolyte and rapid reaction.
The cycling performance of HLRMS and ILRMS cathodes tested at a rate of 1 C in the potential range of 2.0-4.8 V is shown in Fig. 5b. It can be seen that the discharge capacity of
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HLRMS has almost not declined after 100 cycles at 1 C. After 200 cycles, discharge capacities of 174.5 mAh·g-1 and 111.6 mAh·g-1 were retain respectively for HLRMS and ILRMS, which is 94.7 % and 58.9 % of the first discharge capacity. These results reflect the better stability of
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HLRMS cathode over cycling.
The battery based on HLRMS was cycled at different rates ranging from 0.2 C to 5 C (Fig. 5c). The HLRMS cathode delivers a discharge capacity of 243.7, 206.6, 184.2, 165.7, 138.1
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mAh·g-1 at rate of 0.2 C, 0.5 C, 1 C, 2 C and 5 C respectively. We also compared the rate performance of the HLRMS with that of ILRMS at different rates. The battery based on ILRMS
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was examined under the same conditions. The ILRMS delivers a discharge capacity of 227.1, 184.6, 162.0, 134.3, 93.8 mAh·g-1 at rate of 0.2 C, 0.5 C, 1 C, 2 C and 5 C. We can conclude that the HLRMS shows an obvious improvement on the rate performance. In order to illustrate the advantages of the hollow micro-sphere particles, the morphologies of
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HLRMS and ILRMS after cycling are presented (Fig. 6). These two samples before cycling are both sphere particles with diameter about 6 µm (Fig. 3). After cycling, the spherical particles of HLRMS still can be identified on the electrode (Fig. 6a, b), which indicates the morphology of
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the material barely changed and the structure of HLRMS is stable, while the original morphology of ILRMS can’t be found as indicated in Fig. 6c and d. The hollow structure releases the stress of
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Li+ intercalation/ deintercalation and buffers the volume change of material, thus avoiding the cracking of particles. The destruction of particles leads to increase of contact interface between electrode material and electrolyte. The side reactions of active material and electrolyte result in capacity fading and voltage decay.[36] We can see from the Fig. S2b that the material turned into spinel phase and rock-salt phase in the outer surface, only small fraction of layered phase remained. Spacings of the lattice are ~0.47 nm which approximate to the (003) interplanar
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spacing of the Li-rich layered material. And the transformation from layered phase to spinel phase can also be seen from the exposed original inner region (Fig. S2c and d). This illustrates that the original inner region is exposed and transform into spinel-like phase, which causes the
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deepening of voltage decay.
The visual representation of the discharge voltage drop during the cycle is given in Fig. 7b and 7c. After the discharge capacity normalized, the discharge curves of the materials change in
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different cycles could be seen more clearly. It could be observed from the comparison that the voltage fading of HLRMS is much less than that of ILRMS. The average voltage of the HLRMS
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is 3.49 V at the rate of 1 C. The average voltage retention of HLRMS is 97.0 % after 100 cycles and 93.5 % after 200 cycles (Fig. 7a). The average voltage of ILRMS is 3.48 V. The average voltage retention of ILRMS is 91.3 % after 100 cycles and 86.3 % after 200 cycles. Fig. 7(d, e, f) shows the XRD patterns of the material before and after 100 cycles at 1 C. If the material has
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spinel-like phase, we could observe the peak of the hetero phase before the (104) diffraction peak of the layered phase. It can be seen from the (104) diffraction peak in the XRD pattern that a new diffraction peak appears at about 2θ = 44.2° after charge and discharge of the material, which
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belongs to the spinel phase (440) plane, indicating that transformation of layered phase to spinellike phase occurs after cycling [37]. The spinel phase peak (Isp(440)) of the ILRMS is found to be
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stronger than that of the HRLMS, it means that the spinel-like phase transition of the ILRMS material is more severe than that of the HLRMS after charging/discharging. The charge-discharge process of the Li-rich material generally includes the reactions shown below (E1-E6). The remove of Li+ ions in the layered phase during the charge process is accompanied by the oxidation of the transition metal ions (E1).[38, 39] When the voltage reached about 4.5 V, the activation of Li2MnO3 takes place (E2), and Li+ and O2 are removed to
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form electrochemical active layered MnO2. The process is irreversible and is closely related to the characteristics of high capacity and voltage fading.[40] As shown in the differential capacity curve of the material in Fig. 8, reduction of Co ions occurrs at about 4.2 V (E3), reduction of Ni
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ions happens at about 3.7 V (E4),[41] the inserting of Li+ ions into MnO2 takes place at about 3.3 V (E5) and Li intercalation into the spinel-like phase is at about 3.0 V (E6).[42] As shown in Fig. 8b, the capacity at about 3.0 V increases continuously and the capacity above 3.5 V greatly
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reduces, which implies that the capacity could be assigned to the layered phase decreases and some layered phase transforming into spinel-like phase. The capacity about 3.0 V of HLRMS
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(Fig. 8a) barely increases after 20 cycles and the capacity above 3.5 V is more stable than that of ILRMS (Fig. 8b). Charge: LiMO2→Li++MO2
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Li2MnO3→Li++MnO2+O2 Discharge: Co4+→Co3+ Ni4+→Ni2+
E1 E2
E3 E4 E5
MnO2+xLi+→LixMnO2→LiMn2O4 (spinel)→Li2Mn34+O7(spinel)+Mn2+O
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MnO2+Li+→LiMn3+O2
The spinel phase transformation of Lithium-rich layered cathode materials contains the following specific processes. In charging process, metal ions (Co3+ and Ni2+) are oxidized to higher oxidation valences. With some Ni ions entering the Li layer, the Li-rich spinel Li(Li1/3Ni1/2Mn7/6)O4 structure was mainly formed on the surface region [14]. In discharging process, Ni ions at the TM layer are reduced from higher oxidation valences to +2 valence, but it
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is difficult to restore the Ni ions to the original state [30]. When Li ions are intercalated into MnO2, Mn ions are reduced to +3 valence, while the Mn ions of layered phase remain unchanged (+4 valence). Then the LixMnO2 tended to transform into spinel-like phase LiMn2O4, which in
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turn formed Li2Mn34+O7 and MnO (E6) further by disproportionation. Consequently, the MnO turns into fragments and stay in the electrolyte, resulting in capacity loss of the material (~3.3 V) [43]. Because of these reactions, the spinel phase transformation can be indirectly detected with
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XPS.
Fig. 9 shows XPS diagrams of two materials (HLRMS and ILRMS) before and after charge
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and discharge. In Li1.2Ni0.13Co0.13Mn0.54O2, Ni ions are of +2 valence (~855 eV), and Mn ions are of +4 valence (~642.6 eV). After charge and discharge of HLRMS, Ni 2p3/2 major peak will slightly move to +3 valence (~857.6 eV), and Mn elements of +3 valence (~641.1 eV) and of +4 valence (~642.6 eV) coexist, showing that the valence state of the elements agrees with that
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mentioned above. A new Ni 2p3/2 peak centered at ~849.6 eV is caused by the Ni ions at the Li layer [44]. After charge and discharge of ILRMS, the fraction of Ni3+ ions is much higher than that of HLRMS, and the peak of 849.6 eV is stronger, indicating that more Ni ions of ILRMS
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enter the Li layer on the surface to generate spinel-like phase. In addition, Mn elements of both +3 valence and of +4 valence exist. However, comparing with HLRMS, the content of Mn4+ ions
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is relatively high, and that of Mn3+ elements is rather lower, which agrees with the spinel phase transformation and disproportionated reaction of E6, which indicates the spinel phase transformation of ILRMS is more intense. Meanwhile, it can be seen that the displacement of Ni 2p3/2 major peak of ILRMS towards high valence is significant, which tallies with the fact that the capacity fade of the material is significant. Besides, the Ni elements of high valence on material surface are liable to have side reaction with electrolyte to further deteriorate the
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electrochemical performance of the material. The suppression of phase transformation for HLRMS is attributed to two points: (1) stable secondary particles. The secondary particle may collapse after long term cycles, which lead to the increase of contact surface between electrode
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and electrolyte. The side reaction of active material and electrolyte results in the electrode polarization associate with the voltage fading during charge and discharge. HLRMS can keep its morphology after charge and discharge, which own to cavity with appropriate size. (2) The
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phenomenon of the transformation from layered phase to spinel phase involve microstress during Li intercalation/deintercalation. Stress of Li migration will engendered the migration of transition
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metal ions into the Li layers and the formation of spinel phase [37, 45, 46]. The stress caused by Li+ intercalation/deintercalation can be released to avoid the structure destruction and the spinel phase formation. 4. Conclusions
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Hollow spherical layered Li-rich cathode material Li1.2Ni0.13Co0.13Mn0.54O2 was successfully synthesized via a facile lithiation method by using MnO2 hollow spheres as template. The cathode material HLRMS exhibits a favorable enhancement of cycle stability compared with the
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control group ILRMS, especially in the suppression of voltage fading. The capacity of the
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cathode material reaches up to 94.7 % after 200 cycles at 1 C and the average discharge voltage remains 97.0% after 100 cycles and 93.5% after 200 cycles at 1 C. The improvements of the electrochemical performance can be attributed to the hollow secondary particle structure, which release the stress caused by Li+ intercalation/deintercalation. Thus, the structure stable is guaranteed and transformation from layered phase to spinel phase is suppressed. The demonstration of the Li-rich cathode material with hollow spherical particles could offer an efficient way to suppress the voltage decay and to promote its practical applications.
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Acknowledgments This work was supported by the National 973 Program (2015CB251102), the Key Project of NSFC (U1305246, 21673196, 21621091), and the Fundamental Research Funds for the Central
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Universities (20720150042). References
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Table 1 The lattice and structural parameters of HLRMS and ILRMS. c(Å)
c/a
I(003)/I(104)
HLRMS
2.8478
14.21321
4.9909
2.320
ILRMS
2.84944
14.19371
4.9812
2.214
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a(Å)
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All Figure captions
Fig. 1 SEM images of (a, b) the hollow MnO2 precursor and (c, d) the infarctate
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MnO2 precursor.
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Fig. 2 XRD patterns of HLRMS and ILRMS.
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Fig. 3 SEM images of (a, b) HLRMS, (c, d) ILRMS.
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Fig. 4 (a) Elemental ratios measured by EDX and (b, c, d) XPS spectra of HLRMS and ILRMS.
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Fig. 5 Electrochemical performances of HLRMS and ILRMS. (a) the charge and discharge curves at 0.1 C, (b) the cycle performance tested at 1 C, (c) rate performance (1 C=200 mA·g-1).
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Fig. 6 SEM images of (a, b) HLRMS and (c, d) ILRMS after 100 cycles (at 1C).
Fig. 7 (a) Average voltage of HLRMS and ILRMS tested at 1C, (b, c) normalized
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discharge capacity curves of HLMRS and ILRMS, (d, e) XRD patterns of the electrodes (before and after 100 cycles at 1C) and (f) the enlarge XRD pattern between 43-46° (the asterisked peaks represent the diffraction peaks of Al foil).
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Fig. 8 Differential curves (dQ/dV) of (a) HLRMS and (b) ILRMS.
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Fig. 9 XPS patterns of the electrodes (before and after 100 cycles at 1 C).
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Regulation on the secondary particle hollow structure enables Li1.2Ni0.13Co0.13Mn0.54O2 to serve as cathode electrode for high power lithium ion battery showing excellent electrochemical performance.
It proves that the hollow structure of material is really favorable for material to keep the
It demonstrates that the morphology and structure of rich lithium material exert significant impacts on the voltage attenuation of material. The hollow structure can obviously inhibit the voltage, because it can effectively suppress the transformation from layered phase to
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spinel phase.
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original morphology and structure, greatly improving the cycling stability of the material.