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Surface coating effect on thermal properties of delithiated lithium nickel manganese layer oxide
Journal of Power Sources 282 (2015) 511e519
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Journal of Power Sources journal homepage: www.elsevier.com/locate/jpowsour
Surface coating effect on thermal properties of delithiated lithium nickel manganese layer oxide Chang-Heum Jo a, Dae-Hyun Cho a, Jae-won Lee b, Yashiro Hitoshi c, Seung-Taek Myung a, * a b c
Department of Nano Engineering, Sejong University, Seoul 143-747, South Korea Department of Energy Engineering, Dankook University, Cheonan 330-714, South Korea Department of Chemical Engineering, Iwate University, Morioka, Iwate 020-8551, Japan
h i g h l i g h t s Thermal stability of Li0.3[Ni0.7Mn0.3]O2 is improved by silica and silicon phosphate layer. Heating Li0.3[Ni0.7Mn0.3]O2 evolves oxygen from the crystal structure. The oxygen release induces the phase transformation from rhombohedral to cubic spinel phase. The structural change is related to aggressive exothermic decomposition reaction. The exothermic process is retarded to higher temperature by the presence of coating layers.
a r t i c l e i n f o
a b s t r a c t
Article history: Received 22 December 2014 Received in revised form 29 January 2015 Accepted 13 February 2015 Available online 18 February 2015
The thermal stability of electrochemically delithiated bare, silica-coated and silicon phosphate-coated Li0.3[Ni0.7Mn0.3]O2 is studied with in-situ high temperature X-ray diffraction (HT-XRD), thermogravimetric analysis (TGA), transmission electron microscopy (TEM), differential scanning calorimetry (DSC), and time of flight-secondary ion spectroscopy (ToF-SIMS). For the three delithiated materials, gradual phase transformation appears in the crystal structure in the temperature range of 25e600 C: rhombohedral to salt structure via cubic spinel phase, which results from oxygen evolution from the active materials as noticed in TGA. Coating evidently retards the above phase transition toward a high temperature approximately over 40 C owing to less amount of oxygen release from the crystal structure. This effect appears more prominent in the presence of the silicon phosphate coating layer relative to the silica, presumably due to presence of the SiePeO covalent character. Also, the surface layer remains up to 600 C, showing original smooth edges. Therefore, thermal degradation of the active materials is delayed when their surfaces are modified by nanoscale coating layers. © 2015 Elsevier B.V. All rights reserved.
Keywords: Coating Delithiated cathode Thermal stability Lithium Battery
1. Introduction Lithium batteries have drawn significant attention as the most promising energy storage devices [1e3]. Recent research toward the lithium battery usually focuses on large-scale batteries pursuing a long cycle life and high power for vehicles and energy storage applications [4,5]. For these reasons, Ni-rich materials have been intensively studied because of their high energy density and power [6,7]. However, these materials encounter several problems such as
* Corresponding author. E-mail address: [email protected] (S.-T. Myung). http://dx.doi.org/10.1016/j.jpowsour.2015.02.068 0378-7753/© 2015 Elsevier B.V. All rights reserved.
thermal runaway, gas evolution and structural instability [8]. In this regard, demonstration of cell performance under extreme conditions such as high voltage and high temperature is important to understand the mechanism related to the deeply charged electrode materials [9e11]. Especially, the thermal instability of charged materials provides new insight to improve the cathode materials that ensure the safety of batteries [12,13]. These require further studies on the thermal decomposition process and elucidation of the related mechanism. LiNiO2 is a representative cathode material because of its high specific capacity [14e16]. However, the material suffers from structural disorder, in particular cation mixing [17], and thermal instabilities. Hence, the Ni site is usually substituted for other
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metals such as Al, Mn, and Co so as to stabilize the crystal structure. Although Al incorporation into the LiNiO2 structure is effective in improving structural and thermal stabilities, a decrease in the resulting capacity is seen because Al is an electrochemically inactive element (LiNiO2: 175 mAh g1 in 2.5 Ve4.15 V, LiNi0.9Al0.1O2: 125 mAh g1 in 2.5 Ve4.15 V) [18]. In consideration of capacity, therefore, the substitution of Mn or Co into the Ni site is preferred since the trivalent Co or Mn is electrochemically active [19,20]. For example, partial replacement by Mn greatly improves the thermal properties, which is effective in raising the exothermic decomposition temperature of the active materials [21]. In particular, surface modifications with metal oxides [22e24], metal fluorides [25,26], and metal phosphate [27,28] are known to be available to further improve the thermal stability of deeply charged electrodes, namely, the delay of exothermic reaction and the reduction of generated heat [29]. Although there are high capacity materials such as Li[Ni0.8Co0.15Al0.05]O2 and Li[Ni0.8Co0.1Mn0.1]O2, we choose Co-free Li [Ni0.7Mn0.3]O2 material, which delivers a reasonable discharge capacity of about 170 mAh g1, in consideration of material cost. Here, silica and silicon phosphate are employed to modify the surfaces of Li[Ni0.7Mn0.3]O2. Silica is usually used as a filler to reduce thermal deformation of polymer electrolyte [30]. Precipitation of either silica or silicon phosphate nanolayers on the surface of active materials would have positive effect to endow ionic conductivity because of their ionic conducting character [31]. According to our prior reports [32], particle size affects the exothermic reaction; as the surface area becomes larger, a greater amount of heat is generated. To exclude this factor, we select Li[Ni0.7Mn0.3]O2 synthesized via coprecipitation, of which the spherical secondary particles range 9e10 mm in diameter. The presence of divalent Ni in Li layers is found to be approximately 5.9% as previously calculated by the Rietveld refinement [33]. Our effort to reduce cation mixing results in less occupation of divalent Ni in Li layers at approximately 3.9%. In this study, we compare the thermal behavior of the electrochemically delithiated bare Li0.3[Ni0.7Mn0.3]O2 and its modified forms with silica and silicon phosphate coating layers. 2. Experimental Co-free nickel-rich oxides Li[Ni0.7Mn0.3]O2 was synthesized via co-precipitation as follows. An aqueous solution of NiSO4 and MnSO4 with a concentration of 2.0 mol dm3 was pumped into a continuously stirred tank reactor under nitrogen atmosphere. At the same time, a NaOH aqueous solution of 2.0 mol dm3 and the desired amount of NH4OH aqueous solution (a chelating agent) were also separately fed into the reactor. Then, the [Ni0.7Mn0.3](OH)2 particles obtained were filtered, washed, and dried in air. The hydroxide and LiOH powders at a molar ratio 1:1.03 were mixed thoroughly. The mixture was heated at 480 C for 5 h and calcined at 820 C for 15 h in air, then slowly cooled to room temperature at a rate of 1 C min1. The excess Li was used to compensate for the loss of lithium during calcination. To produce silica-coated (hereafter referred to S-coated) Li [Ni0.7Mn0.3]O2 and silicon phosphate-coated (hereafter referred to SP-coated) Li[Ni0.7Mn0.3]O2, TEOS (tetra ethyl orthosilica, Aldrich) and phosphoric acid (Kanto) including TEOS were first separately dissolved completely in anhydrous ethanol at room temperature. The as-synthesized active material, Li[Ni0.7Mn0.3]O2, was slowly poured into the solution. The starting ratio of active material versus coating media was 99.5:0.5 in weight. Then, the solution containing the active material was constantly stirred at 80 C, accompanied by slow evaporation of the solvent. The solution-treated Li[Ni0.7Mn0.3] O2 powders were fired at 600 C for 4 h in air, respectively. For electrochemical delithiation, the prepared active materials
(94 wt%) were mixed with carbon black (3 wt%) and polyvinylidene fluoride (3 wt%) in N-methylpyrrolidinon. The slurry was then cast onto aluminum foil and dried at 110 C for 12 h in a vacuum oven, and then, the electrodes were cut into 3 5 cm dimensions. The electrochemical delithiation was done with aluminum pouch cells using lithium metal as a counter electrode. The assembled cells were charged to 4.3 V by applying a constant current of 50 mA g1 at 25 C. After charging, the cells were disassembled in a glove box, and the electrodes were rinsed salt-free dimethyl carbonate for a day. The electrodes, were then centrifuged in N-methyl-2-pyrrolidone (NMP) to completely remove the polyvinylidene fluoride binder and conducting carbons. After these process, we could obtain a binder-free and carbon-free bare, S-coated, and SP-coated Li0.3[Ni0.7Mn0.3]O2 powders after vacuum drying at 80 C for a day, and they were subjected to various analyses: X-ray diffraction (XRD), thermogravimetric analysis (TGA), high-temperature X-ray diffraction (HT-XRD), transmission electron microscopy (TEM), time-of-flight secondary ion mass spectroscopy (ToF-SIMS), and differential scanning calorimetry (DSC). The crystalline phase of the synthesized products was characterized by powder XRD (Rint-2000, Rigaku) using CuKa radiation. The FULLPROF Rietveld program was used to analyze the powder XRD patterns [34]. The electrochemically delithiated powders were subjected to TGA (loaded sample amount: 10 mg, DTG-60, SHIMADZU, Japan) combined with in-situ HT-XRD (XRD, Rint-2000, Rigaku). The HT-XRD patterns were collected on a Pt heating strip in air. For the TGA, samples were heated from room temperature to 600 C at a heating and a cooling rate of 1 C min1 and were held at selected temperatures for 10 min prior to data collection. The insitu HT-XRD data were obtained at 2q ¼ 10e80 , with a step size of 0.03 and a count time of 0.5 s at each setting temperature. The lattice parameters were calculated by the following method: the positions of the individual peaks were fitted with a pseudo-Voigt or Lorentz function, and typically, peak positions were input to minimize the least-square difference between the calculated and measured peak positions by adjusting the lattice constant and the vertical displacement of the sample. We employed high-resolution transmission electron microscopy (HR-TEM, JEM-3010, JEOL) to observe the phase evolution of the electrochemically delithiated materials. For the TEM measurement, electrochemically delithiated materials were heated at 200e600 C in air for 1 h and quenched to room temperature. To investigate the temperature dependence of each coating layer, the electrochemically delithiated materials were examined by ToF-SIMS (ULVAC-PHI TFS2000, PerkineElmer), equipped with a liquid Biþ ion source and pulse electron flooding and operated at 109 torr. During the analysis, the targets were bombarded by pulsed 10 keV Biþ beams over a 12 12 mm2 area. For the DSC experiment, the electrochemically delithiated Scoated and SP-coated Li[Ni0.7Mn0.3]O2 powders (3e5 mg) were loaded in a stainless-steel sealed pan with a gold-plated copper seal (which can withstand 150 atm of pressure before rupturing and has a capacity of 30 mL). These measurements were carried out in a Pyris 1 calorimeter (PerkineElmer) under a temperature scan rate of 1 C min1. The weight was constant in all cases, indicating that no leaks occurred during the experiments. 3. Result and discussion As-synthesized Li[Ni0.7Mn0.3]O2 is crystallized into typical aNaFeO2 structure with R3 m space group (Fig. 1a). Surface modified SeLi[Ni0.7Mn0.3]O2 and SP-Li[Ni0.7Mn0.3]O2 also show the same crystal structure as that of the bare material (Fig. 1aec). Rietveld refinement of XRD data indicated that the presence of Ni2þ in Li
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Fig. 1. Rietveld refinement result of XRD data for (a) bare, (b) silica-coated and (c) silicon phosphate-coated Li[Ni0.7Mn0.3]O2; electrochemically delithiated (d) bare, (e) silica-coated and (f) silicon phosphate-coated Li0.3[Ni0.7Mn0.3]O2.
layer is approximately 3.9% and that the cation mixing does not vary even after the coating with silica and silicon phosphate. Also, variation in the calculated lattice parameters is negligible after the coating process. Electrochemical delithiation was carried out using these samples to avoid change derived during chemical delithiation, often resulting in morphological degradation [35]. The electrodes are first charged to 4.3 V, delivering a charge capacity of approximately 195 mAh g1, which can be expressed as Li0.3[Ni0.7Mn0.3]O2 (Fig. 2). The delivered discharge capacity reaches 170 mAh g1, which is much improved than that of our prior report, presumably due to minimization of cation mixing in the crystal structure [33]. The charged electrodes were rinsed with salt-free dimethyl carbonate for a week, in an Ar filled glove box. And, the XRD patterns of delithiated sample were calibrated using the internal standard of the Al current collector (Fig. 1def). Similar to the fresh materials, Rietveld refinement of XRD data for these delithiated electrodes are carried out by assuming R3 m space group. The original hexagonal structure is kept after the electrochemical delithiation for all samples. Refined data indicate contraction of the a-axis due to oxidation of transition metals while expansion of the c-axis resulting from increased ionicity in the oxide lattice (Table 1). Occupation of Ni2þ in Li layer also slightly increased in comparison with those of fresh materials, presumably due to migration of Ni
element on the electrochemical delithiation. Fig. 3 shows the surface morphology of the fresh and electrochemically delithiated bare, S- and SP-coated Li[Ni0.7Mn0.3]O2
Fig. 2. First charge curves of bare, silica-coated, silicon phosphate-coated Li [Ni0.7Mn0.3]O2.
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Table 1 Rietveld refinement result of XRD data for bare, silicon-coated and silicon phosphate-coated Lix[Ni0.7Mn0.3]O2 (x ¼ 1, 0.3).
Bare Li[Ni0.7Mn0.3]O2 Silica coated-Li[Ni0.7Mn0.3]O2 Silicon phosphate coated-Li[Ni0.7Mn0.3]O2 Bare Li0.3[Ni0.7Mn0.3]O2 Silica coated-Li0.3[Ni0.7Mn0.3]O2 Silicon phosphate coated-Li0.3[Ni0.7Mn0.3]O2
a-axis/Å
c-axis/Å
Ni2þ in Li layer/%
Rwp/%
2.8831(1) 2.8830(1) 2.8833(1)
14.2568(4) 14.2569(4) 14.2566(3)
3.9(3) 3.9(3) 3.9(2)
12.5 12.4 14.0
2.8381(2) 2.8385(1) 2.8384(1)
14.4035(4) 14.4033(4) 14.4031(2)
5.8(2) 5.8(3) 5.8(3)
10.7 12.9 12.7
particles. The fresh bare particles have smooth surfaces, and there are no additional layers on the surface (Fig. 3a). For S- and SPcoated particles, in contrast, surface layers with thickness about 10 nm are noticed on the surface of the particles (Fig. 3b and c). Even after the electrochemical delithiation up to Li0.3[Ni0.7Mn0.3]O2, the smooth surfaces are still observed for the coated particles (Fig. 3e and f), while the delithiated bare particles seems to be somehow damaged during the delithiation (Fig. 3d), which demonstrates that the presence of coating layer is effective in the protection of the outer surface during migration of Liþ ions. The surfaces for the electrochemically delithiated bare and coated electrodes are analyzed by ToF-SIMS (Fig. 4). In the highlighted ToF-SIMS spectra, a Si2Oþ (m ¼ 71.94) fragment is detected on the coated electrodes whereas the fragment is absent for the bare. Since the silicon phosphate consists of SiePeO bond, the appearance of Si2Oþ fragment is reasonable. And SiPOþ (m ¼ 74.94) fragment is found for the SP-coated electrode (bottom of Fig. 4b). These data coincide with the TEM results that the coating layer remains on the surface of active materials after the electrochemical
delithiation. TGA investigation was carried out for the electrochemically delithiated bare, S-, and SP-coated powders (Fig. 5). The charged powders show distinct differences in weight variation on heating. Since the powders were dried under vacuum at 80 C for a day, the weight loss associated with evaporation of water is negligible up to 150 C for all samples. Then, the first drastic weight loss begins from 200 C and is continued to 400 C, showing 9.7 wt. % of weight loss for the delithiated bare, which results from oxygen release from the parent material [29]. The second weight loss is observed in the temperature range of 400e550 C (13.2 wt. %), after which the weight keeps constant up to 600 C. For SeLi0.3[Ni0.7Mn0.3]O2, the initiation temperature of the drastic weight loss is somehow delayed to 240 C, after which the weight loss is progressed to 460 C with 8.8 wt. % of oxygen loss. The weight further decreases to 600 C approximately 8.5 wt. %. A similar tendency is observed for the SP-Li0.3[Ni0.7Mn0.3]O2. Namely, the initiation temperature of the drastic weight loss is further delayed to 260 C, and the weight loss is progressed relatively slow to 460 C (5.3 wt. %) and then accelerated to 510 C (3.3%), which belongs to the first stage of the weight loss process. In comparison with the weight loss process for S- and SP-Li0.3[Ni0.7Mn0.3]O2, it is obvious that the silicon phosphate coating (5.5 wt. %) seems to be more effective in suppressing weight loss than silica coating to 460 C (8.8 wt. %) because of less evolution of oxygen from the crystal structure of active materials. Further weight loss approximates to 7 wt. % in the temperature range of 510e600 C, which is similar to the silica coating. As a result, it is clear that oxygen release is more dominant and starts relatively at lows temperature for the bare materials, while the presence of a coating layer obviously delays the oxygen release process upon heating. To further correlate the weight loss to the structural change from the charged materials upon heating, in-situ HT-XRD
Fig. 3. TEM images for (a) bare, (b) silica-coated and (c) silicon phosphate-coated Li[Ni0.7Mn0.3]O2; electrochemically delithiated (d) bare, (e) silica-coated and (f) silicon phosphatecoated Li0.3[Ni0.7Mn0.3]O2.
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Fig. 4. ToF-SIMS results of bare, silica-coated and silicon phosphate-coated Lix[Ni0.7Mn0.3]O2 (x ¼ 1 (a, b); x ¼ 0.3 (c, d)). (top: bare, middle: silica-coated material, bottom: silicon phosphate-coated material).
measurements were carried out for the electrochemically delithiated bare, S-, and SP- Li0.3[Ni0.7Mn0.3]O2 (Fig. 6aec). The series of patterns were collected at an interval of 10 C in the temperature range of 150e300 C. Otherwise, the interval was 50 C. Bare, S- and SP-Li0.3[Ni0.7Mn0.3]O2 have a stable rhombohedral structure at room temperature. Although the majority of crystal structure remains as the rhombohedral structure, the resulting peak positions are shifted and the intensity increases or decreases to certain temperatures; namely, for bare to 200 C, for SeLi0.3Ni0.7Mn0.3O2 to 240 C, and SP-Li0.3Ni0.7Mn0.3O2 to 260 C. Also, the relative intensity ratio of I(003)/I(104) shows that the occupation of transition metal ions in the Li layers decreases gradually without formation of new phases in the same temperature range: bare: 1.03, SeLi0.3Ni0.7Mn0.3O2: 1.22, and SP-Li0.3Ni0.7Mn0.3O2: 1.29 at 50 C; bare: 0.85, SeLi0.3Ni0.7Mn0.3O2: 0.94, SP-Li0.3Ni0.7Mn0.3O2: 0.95 at 150 C; and bare: 0.68, SeLi0.3Ni0.7Mn0.3O2: 0.79, and SPLi0.3Ni0.7Mn0.3O2: 0.88 at 200 C. Lowering of the ratio is obvious with increasing the heating temperature, meaning that migration of transition metal, particularly nickel, to the Li layers is accelerated from the octahedral 3b sites to vacant 3a sites via tetrahedral 6c sites in the Li layer, which is face-shared with the 3a and 3b sites,
Fig. 5. TGA curves of bare, silica-coated and silicon phosphate-coated Li0.3[Ni0.7Mn0.3] O2.
with increasing temperature. In this charged state, in which both Ni and Mn are oxidized to tetravalent, it is likely that energetically unstable Ni4þ (0.48 Å) relative to Mn4þ, which has smaller ionic size than that of Mn4þ (0.53 Å), would readily migrate to Li layers. This behavior results in the increment of the a-axis and simultaneously affects the shrinkage of LiO2 slabs accompanying decrease in the caxis. Hence, the gradual decrease in the c/a ratio is observed to 200 C (Fig. 7). This result demonstrates that cation migration is continuously progressed in the rhombohedral symmetry without structural changes to 200 C for the bare, 240 C for the SeLi0.3Ni0.7Mn0.3O2, and 260 C for the SP-Li0.3Ni0.7Mn0.3O2. The advent of the cubic spinel structure is initiated above these temperatures. These onset temperatures are almost consistent with the starting point of weight loss in the TGA result (Fig. 5). In particular, the (108) and (110) peaks indicating a typical O3 type layer structure are merged to the (440) peak of cubic spinel structure with Fd3m space group. Since oxygen evaporation is aggressively progressed in the temperature region of 200e300 C for the bare, the new spinel (311), (400), and (440) peaks evolve fast in this temperature. And the structural transformation to the cubic spinel phase seems to be terminated at 400 C (Fig. 6a), which coincides with the first weight loss region in TGA (Fig. 5). Interestingly, these peaks shift toward lower angle, indicating thermal expansion of the crystal lattice. Above the temperature of 400 C, formation of rock salt type phase is noticed and the two phases, MO (Fm3m) and M3O4 (M: Ni and Mn) coexist up to 600 C. This temperature range belongs to the second oxygen evolution region, approximately 13.2 wt. %. During the oxygen evaporation, structural reorganization occurs via the continuous migration of cations to a more stable site such as the vacant lithium site and tetrahedral site, resulting in the formation of a simple cubic MO compound. The two phases are still observed even after cooling to room temperature. Different from the bare Li0.3Ni0.7Mn0.3O2, the SeLi0.3Ni0.7Mn0.3O2 undergoes a slow phase evolution toward cubic spinel structure in the first weight loss region, 240e460 C (Fig. 6b). At 250 C, the rhombohedral structure begins to be transformed to the cubic spinel structure, and these two phases coexist to 400 C. There are apparently no shifts of diffraction peaks to this temperature, probably due to the rearrangement of crystal structure to cubic spinel structure. As soon as the single cubic spinel phase is formed at 450 C, the diffraction peak positions move toward the
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Fig. 6. In-situ HT-XRD patterns of (a) bare, (b) silica-coated and (c) silicon phosphate-coated Li0.3[Ni0.7Mn0.3]O2 in the temperature range of 25 Ce600 C.
lower angle owing to the thermal expansion of the structure. The temperature is close to the ending point of the first weight loss in TGA (Fig. 5). The released oxygen content is about 8.8 wt. % to 460 C for the SeLi0.3Ni0.7Mn0.3O2, while the content is approximately 9.7 wt. % to 400 C for the bare Li0.3Ni0.7Mn0.3O2. Above 450 C, the advent of the simple cubic MO is observed (Fm3m) and M3O4 (M: Ni and Mn) coexist up to 600 C. In the temperature range of 450e600 C, the evolved amount of oxygen from the parent material is about 8.5 wt. %, which is lower than that of bare. Therefore, the phase transformation from the cubic spinel to the simple cubic phase is delayed. The two phases are still observed even after cooling to room temperature. These data suggest that the presence of silica coating layers suppresses oxygen evolution from the crystal structure, and this, in turn, leads the phase transition tedious. For the SP-Li0.3Ni0.7Mn0.3O2, the phase transformation toward the cubic spinel phase is observed at a slightly higher temperature (270 C, Fig. 6c) than that of the SeLi0.3Ni0.7Mn0.3O2 (250 C, Fig. 6b). The cubic spinel structure is still perceived up to 450 C due to the less amount of oxygen evolution to 510 C (5.5 wt. %), which
confirms the effect of coating that is able to delay the phase transformation triggered by the oxygen release from the parent oxide. Above 500 C, the resulting phase is a mixture of the cubic spinel and simple cubic phases. It is likely that the silicon phosphate layer is favored rather than silica layer because the series of phase transformation is delayed to higher temperature. Therefore, the phase transformation is initiated by the gradual oxygen release, which simultaneously causes rearrangement of structure via migration of metal elements, from the structure in the following sequence; (i) rhombohedral phase, (ii) coexistence rhombohedral and spinel phases, (iii) spinel phase, and (iv) coexistence of spinel and rock salt phases. As a result, the consecutive cation migration resulting from oxygen release gives rise to variation in lattice parameters (Fig. 7a). For convenience, the lattice parameters are represented as the rhombohedral symmetry even after the phase transformation to the cubic spinel structure. Drastic weight change reflects an abrupt decrease in the c/a ratio because of the advent of cubic spinel structure in the temperature of 200e450 C. Among samples, the bare material undergoes a fast structural transformation,
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Fig. 7. Variation in of c/a ratio for the electrochemically delithiated bare, silica-coated and silicon phosphate-coated Li0.3[Ni0.7Mn0.3]O2 in the temperature range of 25 Ce600 C.
terminating at about 300 C, which shows the c/a ratio close to 4.9 that is interpreted as the cubic symmetry. In comparison with the prior reports, the onset temperature of phase transformation from rhombohedral to the spinel phase for the present Li0.3[Ni0.7Mn0.3] O2 (220 C) was similar to that of Li0.3[Ni0.7Co0.15Al0.15]O2 (230 C) [18] and Li0.33[Ni0.5Co0.2Mn0.3]O2 (223 C) apart from the different ratio of cation mixing in the Li layers [35,36]. By contrast, the variation is retarded by the presence of foreign silica and silicon phosphate coating layers. Our prior report revealed the effect of AlF3 coating on Li0.35[Ni1/3Co1/3Mn1/3]O2 such that the coating layer blocks oxygen evolution during exothermic decomposition, resulting in delayed phase transition [29]. It is likely that the present silica and silicon phosphate layers play the similar role to the AlF3 layer, since the TGA and in-situ HT-XRD results are similar to the previous results even in Ni-rich Li0.3Ni0.7Mn0.3O2 composition. The electrochemically delithiated electrodes were heated at 250 C, 400 C and 600 C, whose temperatures correspond to phase transition points, in air for 1 h and quenched to room temperature for surface and ToF-SIMS analyses. Fig. 8 exhibits the surface morphology and the chemistry information for the bare, S-, and SP-Li0.3Ni0.7Mn0.3O2 powders quenched at 250 C. For the bare
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material, outer part is damaged, and it progresses toward the particle inside presumably due to the oxygen evolution from the oxide particle after the quenching at 250 C (Fig. 8a). The appearing cubic spinel phase in the in-situ XRD patterns would be ascribed to the damaged surface after the oxygen evolution. Unlike the bare particle, the S- and SP-Li0.3Ni0.7Mn0.3O2 particles present smooth and clear surfaces (Fig. 8b and d). In addition, the silica and silicon phosphate coating layers are still visible, as further evidenced from the Si2Oþ fragment and silicon phosphate layer from the SiPOþ fragment (Fig. 8c and e), indicating the presence of coating layers without thermal decomposition at this temperature. Oxygen would be aggressively removed from the oxide particle for the bare material at 250 C as evident from TGA, whereas the oxygen evolution is just initiated at the temperature for the coated samples. At 400 C (Fig. 9a), better particle stability is noticed for both coated particles (Fig. 9b and d), and the coating layers are still observed. It seems that the coated particles remained intact, although there was weight loss of approximately 6.7 wt. % for the SeLi0.3Ni0.7Mn0.3O2 and 3.7 wt. % for the SP-Li0.3Ni0.7Mn0.3O2. Less deterioration observed in those particles is reasonable because the quenching was done at the intermediate temperature of the phase transformation. Additional information is given by the ToF-SIMS spectra that the silica, as the SiOþ fragment, and silicon phosphate, as the SiPOþ fragment, coating layers are not decomposed but still appear on the outermost surface (Fig. 9c and e). At 600 C, the crystal structures of the three delithiated materials are stabilized into the simple cubic structure as a result of the two-step oxygen evolution from the parent particles. Unlike the bare material that starts to form the rock salt phase from 400 C, the advent of the rock salt phase is delayed for the coated materials: 450 C for the SeLi0.3Ni0.7Mn0.3O2 and 500 C for the SPLi0.3Ni0.7Mn0.3O2, which would result from the less damage of the delithiated active materials even at 600 C. In addition, coating media remained on the outermost surface of the active material as the presence of coating layers are clearly observed on the TEM images and ToF-SIMS spectra (Fig. 10bee). Note that the coating layers are very thin, below 10 nm (Fig. 3). The coating layers probably block the oxygen evolution from the oxide lattice. As evident from the TGA, XRD, TEM, and ToF-SIMS results, the SP-Li0.3Ni0.7Mn0.3O2 experiences slow phase
Fig. 8. TEM images and ToF-SIMS results of the bare (a), silica coated (b, c) and silicon phosphate coated Li0.3[Ni0.7Mn0.3]O2 (d, e) quenched at 250 C. (in ToF-SIMS result, upper: Si2Oþ fragment, bottom: SiPOþ fragment).
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Fig. 9. TEM images and ToF-SIMS result of the bare (a), silica coated (b, c) and silicon phosphate coated Li0.3[Ni0.7Mn0.3]O2 (d, e) quenched at 400 C. (in ToF-SIMS result, upper: Si2Oþ fragment, bottom: SiPOþ fragment).
transformation due to less evaporation of oxygen from the parent oxide. It is thought that the presence of PeO covalent bond may assist to form the coating layer much solid relative to the silica coating, and this, in turn, would further delay the oxygen evaporation from the delithiated Li0.3Ni0.7Mn0.3O2. Fig. 11 represents DSC data for the electrochemically delithiated bare, S-, and SP-Li0.3Ni0.7Mn0.3O2 in the presence of electrolyte. The onset temperatures of exothermic reaction coincide with the TGA and in-situ HT-XRD data for the three samples, in which the cubic spinel phase begins to be produced as a result of oxygen loss from the rhombohedral oxide lattice. The main exothermic decomposition occurs at a somehow elevated temperature by the silica and silicon phosphate coating, and the generated heat is also dependent on the coating layer. Specifically, less evolution of oxygen brings about less exothermic heat as follows: bare (~614.4 J g1), silica coated-material (~498.6 J g1), and silicon phosphate-coated
material (~432.7 J g1). It is most likely that the coating layers block the oxygen release from the active material, and it, in turn, affects the shift of the main exothermic reaction to a higher temperature. After the reaction, the majority of the resulting phase is composed of the cubic spinel phase as confirmed in the in-situ HTXRD. Therefore, the presence of coating layers on the surface of the delithiated active materials is responsible for the improved thermal properties. 4. Conclusion The thermal behavior of the delithiated bare, silica-coated and silicon phosphate-coated Li0.3[Ni0.7Mn0.3]O2 is investigated. Oxygen release from the parent oxide materials induces phase transformation in a sequence of rhombohedral to cubic spinel and finally the simple cubic phase. Applying the coating layers on the
Fig. 10. TEM images and ToF-SIMS result of the bare (a), silica coated (b, c) and silicon phosphate coated Li0.3[Ni0.7Mn0.3]O2 (d, e) quenched at 600 C. (in ToF-SIMS result, upper: Si2Oþ fragment, bottom: SiPOþ fragment).
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References
Fig. 11. Comparison of DSC traces of the bare, silica coated and silicon phosphate coated Li0.3[Ni0.7Mn0.3]O2.
surface of cathode materials, the onset temperature of the exothermic decomposition is retard over 40 C compared to the bare material because the oxygen loss is suppressed by the coating layers. In particular, silicon phosphate coating is effective in delaying the exothermic reaction relative to the oxide coating due to the sturdiness of the coating layer supported by the presence of covalent bond of silicon phosphate. Therefore, the role of the coating layer and the related mechanism behind the thermal properties are comprehensively understood by means of TGA, TEM, in-situ HT-XRD, and ToF-SIMS, which are used to interpret the DSC behavior. Acknowledgments The authors thank Miwa Watanabe, Iwate University, for her helpful assistance in the experimental work. This research was supported by Basic Science Research Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Education, Science and Technology (NRF-2009-0093467). This work was also supported by MKE/KEIT (10041856 and 10041094). This research was supported by Basic Science Research Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Education, Science and Technology (2014R1A2A1A11051197)
[1] B. Scrosati, J. Garche, J. Power Sources 195 (2010) 2419. [2] M. Winter, J.O. Besenhard, M.E. Spahr, P. Novak, Adv. Mater. 10 (1998) 725. [3] V. Etacheri, R. Marom, R. Elazari, G. Salitra, D. Aurbach, Energy Environ. Sci. 4 (2011) 3243. [4] L. Lu, X. Han, J. Li, J. Hua, M. Ouyang, J. Power Sources 226 (2013) 272. [5] M.M. Thackeray, C. Wolverton, E.D. Issacs, Energy Environ. Sci. 5 (2015) 7854. [6] Y.-K. Sun, Z. Chen, H.-J. Noh, D.-J. Lee, H.-G. Jung, Y. Ren, S. Wang, C.-S. Yoon, S.-T. Myung, K. Amine, Nat. Mater. 11 (2012) 942. [7] P.Y. Liao, J.G. Duh, S.R. Sheen, J. Electrochem. Soc. 152 (2005) A1695. [8] A. Wuersig, W. Scheifele, P. Nov ak, J. Electrochem. Soc. 154 (2007) A449. [9] X. Feng, J. Sun, M. Ouyang, X. He, L. Lu, X. Han, M. Fang, H. Peng, J. Power Sources 272 (2014) 457. [10] J. Lamb, C.J. Orendorff, K. Amine, G. Krumdick, Z. Zhang, L. Zhang, A.S. Gozdz, J. Power Sources 247 (2014) 1011. [11] Z. Chen, Y. Ren, E.J. Lee, C. Johnson, Y. Qin, K. Amine, Adv. Energy Mater. 6 (2013) 729. [12] L. Wu, K.W. Nam, X. Wang, Y. Zhou, J.C. Zheng, X.Q. Yang, Y. Zhu, Chem. Mater. 23 (2011) 3953. [13] K.W. Nam, S.M. Bak, E. Hu, Z. Yu, Y. Zhou, X. Wang, L. Wu, Y. Zhu, K.Y. Chung, X.Q. Yang, Adv. Funct. Mater. 23 (2013) 1047. [14] J. Molenda, P. Wilk, J. Marzec, Solid State Ionics 146 (2002) 73. [15] B.H. Kim, J.H. Kim, I.H. Kwon, M.Y. Song, Ceram. Int. 33 (2007) 837. [16] Z. Liu, H. Zhen, Y. Kim, C. Liang, J. Power Sources 196 (2011) 10201. [17] N. Yabuuchi, Y.T. Kim, H.H. Li, S.H. Yang, Chem. Mater. 20 (2008) 4936. [18] M. Guilmard, A. Rougier, A. Grune, L. Croguennec, C. Delmas, J. Power Sources 115 (2003) 305. [19] A. Rougier, I. Saadoune, P. Gravereau, P. Willmann, C. Delmas, Solid State Ionics 90 (1996) 83. [20] S.T. Myung, S. Komaba, K. Kurihara, K. Hosoya, N. Kumagai, Y.K. Sun, I. Nakai, M. Yonemura, T. Kamiyama, Chem. Mater. 18 (2006) 1658. [21] J. Jiang, K.W. Eberman, L.J. Krause, J.R. Dahn, J. Electrochem. Soc. 152 (2005) A566. [22] S.T. Myung, K. Izumi, S. Komaba, Y.K. Sun, H. Yashiro, N. Kumagai, Chem. Mater. 17 (2005) 2427. [23] S.T. Myung, K. Izumi, S. Komaba, H. Yashiro, H.J. Bang, Y.K. Sun, N. Kumagai, J. Phys. Chem. C 111 (2007) 4061. [24] M. Bettge, Y. Li, B. Sankaran, N.D. Rago, T. Spila, R.T. Haasch, I. Petrov, D.P. Abraham, J. Power Sources 233 (2013) 346. [25] K.S. Lee, S.T. Myung, K. Amine, H. Yashiro, Y.K. Sun, J. Mater. Chem. 19 (2009) 1995. [26] K.J. Rosina, M. Jiang, D.L. Zeng, E. Salager, A.S. Best, C.P. Grey, J. Mater. Chem. 22 (2012) 20602. [27] A.T. Appapillai, A.N. Mansour, J.-P. Cho, Y. Shao-Horn, Chem. Mater. 19 (2007) 5748. [28] D.J. Lee, B. Scrosati, Y.K. Sun, J. Power Sources 196 (2011) 7745. [29] S.T. Myung, K.S. Lee, C.S. Yoon, Y.K. Sun, K. Amine, H. Yashiro, J. Phys. Chem. C 114 (2010) 4710. [30] K.H. Teoh, S. Ramesh, A.K. Arof, J. Solid State Electrochem. 16 (2012) 3165. [31] D. Praveen, S.V. Bhat, R. Damle, Ionics 19 (2013) 1375. [32] K.S. Lee, S.T. Myung, J.S. Moon, Y.K. Sun, J. Electrochem. Soc. 161 (2014) A920. [33] D.H. Cho, C.H. Jo, W.S. Cho, Y.J. Kim, H. Yashiro, Y.K. Sun, S.T. Myung, J. Electrochem. Soc. 161 (2014) A920. [34] T. Roisnel, J. Rodriguez-Carjaval, Fullprof Manual, Institut Laue-Langevin, Grenoble, France, 2002. [35] Y.H. Cho, D.H. Jang, J.B. Yoon, H.C. Kim, T.K. Ahn, K.W. Nam, Y.E. Sung, W.S. Kim, Y.S. Lee, X.Q. Yang, W.S. Yoon, J. Alloys Compd. 562 (2013) 219. [36] M. Guilmard, L. Croguennec, C. Delmas, Chem. Mater. 15 (2003) 4484.
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Surface coating effect on thermal properties of delithiated lithium nickel manganese layer oxide Chang-Heum Jo a, Dae-Hyun Cho a, Jae-won Lee b, Yashiro Hitoshi c, Seung-Taek Myung a, * a b c
Department of Nano Engineering, Sejong University, Seoul 143-747, South Korea Department of Energy Engineering, Dankook University, Cheonan 330-714, South Korea Department of Chemical Engineering, Iwate University, Morioka, Iwate 020-8551, Japan
h i g h l i g h t s Thermal stability of Li0.3[Ni0.7Mn0.3]O2 is improved by silica and silicon phosphate layer. Heating Li0.3[Ni0.7Mn0.3]O2 evolves oxygen from the crystal structure. The oxygen release induces the phase transformation from rhombohedral to cubic spinel phase. The structural change is related to aggressive exothermic decomposition reaction. The exothermic process is retarded to higher temperature by the presence of coating layers.
a r t i c l e i n f o
a b s t r a c t
Article history: Received 22 December 2014 Received in revised form 29 January 2015 Accepted 13 February 2015 Available online 18 February 2015
The thermal stability of electrochemically delithiated bare, silica-coated and silicon phosphate-coated Li0.3[Ni0.7Mn0.3]O2 is studied with in-situ high temperature X-ray diffraction (HT-XRD), thermogravimetric analysis (TGA), transmission electron microscopy (TEM), differential scanning calorimetry (DSC), and time of flight-secondary ion spectroscopy (ToF-SIMS). For the three delithiated materials, gradual phase transformation appears in the crystal structure in the temperature range of 25e600 C: rhombohedral to salt structure via cubic spinel phase, which results from oxygen evolution from the active materials as noticed in TGA. Coating evidently retards the above phase transition toward a high temperature approximately over 40 C owing to less amount of oxygen release from the crystal structure. This effect appears more prominent in the presence of the silicon phosphate coating layer relative to the silica, presumably due to presence of the SiePeO covalent character. Also, the surface layer remains up to 600 C, showing original smooth edges. Therefore, thermal degradation of the active materials is delayed when their surfaces are modified by nanoscale coating layers. © 2015 Elsevier B.V. All rights reserved.
Keywords: Coating Delithiated cathode Thermal stability Lithium Battery
1. Introduction Lithium batteries have drawn significant attention as the most promising energy storage devices [1e3]. Recent research toward the lithium battery usually focuses on large-scale batteries pursuing a long cycle life and high power for vehicles and energy storage applications [4,5]. For these reasons, Ni-rich materials have been intensively studied because of their high energy density and power [6,7]. However, these materials encounter several problems such as
* Corresponding author. E-mail address: [email protected] (S.-T. Myung). http://dx.doi.org/10.1016/j.jpowsour.2015.02.068 0378-7753/© 2015 Elsevier B.V. All rights reserved.
thermal runaway, gas evolution and structural instability [8]. In this regard, demonstration of cell performance under extreme conditions such as high voltage and high temperature is important to understand the mechanism related to the deeply charged electrode materials [9e11]. Especially, the thermal instability of charged materials provides new insight to improve the cathode materials that ensure the safety of batteries [12,13]. These require further studies on the thermal decomposition process and elucidation of the related mechanism. LiNiO2 is a representative cathode material because of its high specific capacity [14e16]. However, the material suffers from structural disorder, in particular cation mixing [17], and thermal instabilities. Hence, the Ni site is usually substituted for other
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metals such as Al, Mn, and Co so as to stabilize the crystal structure. Although Al incorporation into the LiNiO2 structure is effective in improving structural and thermal stabilities, a decrease in the resulting capacity is seen because Al is an electrochemically inactive element (LiNiO2: 175 mAh g1 in 2.5 Ve4.15 V, LiNi0.9Al0.1O2: 125 mAh g1 in 2.5 Ve4.15 V) [18]. In consideration of capacity, therefore, the substitution of Mn or Co into the Ni site is preferred since the trivalent Co or Mn is electrochemically active [19,20]. For example, partial replacement by Mn greatly improves the thermal properties, which is effective in raising the exothermic decomposition temperature of the active materials [21]. In particular, surface modifications with metal oxides [22e24], metal fluorides [25,26], and metal phosphate [27,28] are known to be available to further improve the thermal stability of deeply charged electrodes, namely, the delay of exothermic reaction and the reduction of generated heat [29]. Although there are high capacity materials such as Li[Ni0.8Co0.15Al0.05]O2 and Li[Ni0.8Co0.1Mn0.1]O2, we choose Co-free Li [Ni0.7Mn0.3]O2 material, which delivers a reasonable discharge capacity of about 170 mAh g1, in consideration of material cost. Here, silica and silicon phosphate are employed to modify the surfaces of Li[Ni0.7Mn0.3]O2. Silica is usually used as a filler to reduce thermal deformation of polymer electrolyte [30]. Precipitation of either silica or silicon phosphate nanolayers on the surface of active materials would have positive effect to endow ionic conductivity because of their ionic conducting character [31]. According to our prior reports [32], particle size affects the exothermic reaction; as the surface area becomes larger, a greater amount of heat is generated. To exclude this factor, we select Li[Ni0.7Mn0.3]O2 synthesized via coprecipitation, of which the spherical secondary particles range 9e10 mm in diameter. The presence of divalent Ni in Li layers is found to be approximately 5.9% as previously calculated by the Rietveld refinement [33]. Our effort to reduce cation mixing results in less occupation of divalent Ni in Li layers at approximately 3.9%. In this study, we compare the thermal behavior of the electrochemically delithiated bare Li0.3[Ni0.7Mn0.3]O2 and its modified forms with silica and silicon phosphate coating layers. 2. Experimental Co-free nickel-rich oxides Li[Ni0.7Mn0.3]O2 was synthesized via co-precipitation as follows. An aqueous solution of NiSO4 and MnSO4 with a concentration of 2.0 mol dm3 was pumped into a continuously stirred tank reactor under nitrogen atmosphere. At the same time, a NaOH aqueous solution of 2.0 mol dm3 and the desired amount of NH4OH aqueous solution (a chelating agent) were also separately fed into the reactor. Then, the [Ni0.7Mn0.3](OH)2 particles obtained were filtered, washed, and dried in air. The hydroxide and LiOH powders at a molar ratio 1:1.03 were mixed thoroughly. The mixture was heated at 480 C for 5 h and calcined at 820 C for 15 h in air, then slowly cooled to room temperature at a rate of 1 C min1. The excess Li was used to compensate for the loss of lithium during calcination. To produce silica-coated (hereafter referred to S-coated) Li [Ni0.7Mn0.3]O2 and silicon phosphate-coated (hereafter referred to SP-coated) Li[Ni0.7Mn0.3]O2, TEOS (tetra ethyl orthosilica, Aldrich) and phosphoric acid (Kanto) including TEOS were first separately dissolved completely in anhydrous ethanol at room temperature. The as-synthesized active material, Li[Ni0.7Mn0.3]O2, was slowly poured into the solution. The starting ratio of active material versus coating media was 99.5:0.5 in weight. Then, the solution containing the active material was constantly stirred at 80 C, accompanied by slow evaporation of the solvent. The solution-treated Li[Ni0.7Mn0.3] O2 powders were fired at 600 C for 4 h in air, respectively. For electrochemical delithiation, the prepared active materials
(94 wt%) were mixed with carbon black (3 wt%) and polyvinylidene fluoride (3 wt%) in N-methylpyrrolidinon. The slurry was then cast onto aluminum foil and dried at 110 C for 12 h in a vacuum oven, and then, the electrodes were cut into 3 5 cm dimensions. The electrochemical delithiation was done with aluminum pouch cells using lithium metal as a counter electrode. The assembled cells were charged to 4.3 V by applying a constant current of 50 mA g1 at 25 C. After charging, the cells were disassembled in a glove box, and the electrodes were rinsed salt-free dimethyl carbonate for a day. The electrodes, were then centrifuged in N-methyl-2-pyrrolidone (NMP) to completely remove the polyvinylidene fluoride binder and conducting carbons. After these process, we could obtain a binder-free and carbon-free bare, S-coated, and SP-coated Li0.3[Ni0.7Mn0.3]O2 powders after vacuum drying at 80 C for a day, and they were subjected to various analyses: X-ray diffraction (XRD), thermogravimetric analysis (TGA), high-temperature X-ray diffraction (HT-XRD), transmission electron microscopy (TEM), time-of-flight secondary ion mass spectroscopy (ToF-SIMS), and differential scanning calorimetry (DSC). The crystalline phase of the synthesized products was characterized by powder XRD (Rint-2000, Rigaku) using CuKa radiation. The FULLPROF Rietveld program was used to analyze the powder XRD patterns [34]. The electrochemically delithiated powders were subjected to TGA (loaded sample amount: 10 mg, DTG-60, SHIMADZU, Japan) combined with in-situ HT-XRD (XRD, Rint-2000, Rigaku). The HT-XRD patterns were collected on a Pt heating strip in air. For the TGA, samples were heated from room temperature to 600 C at a heating and a cooling rate of 1 C min1 and were held at selected temperatures for 10 min prior to data collection. The insitu HT-XRD data were obtained at 2q ¼ 10e80 , with a step size of 0.03 and a count time of 0.5 s at each setting temperature. The lattice parameters were calculated by the following method: the positions of the individual peaks were fitted with a pseudo-Voigt or Lorentz function, and typically, peak positions were input to minimize the least-square difference between the calculated and measured peak positions by adjusting the lattice constant and the vertical displacement of the sample. We employed high-resolution transmission electron microscopy (HR-TEM, JEM-3010, JEOL) to observe the phase evolution of the electrochemically delithiated materials. For the TEM measurement, electrochemically delithiated materials were heated at 200e600 C in air for 1 h and quenched to room temperature. To investigate the temperature dependence of each coating layer, the electrochemically delithiated materials were examined by ToF-SIMS (ULVAC-PHI TFS2000, PerkineElmer), equipped with a liquid Biþ ion source and pulse electron flooding and operated at 109 torr. During the analysis, the targets were bombarded by pulsed 10 keV Biþ beams over a 12 12 mm2 area. For the DSC experiment, the electrochemically delithiated Scoated and SP-coated Li[Ni0.7Mn0.3]O2 powders (3e5 mg) were loaded in a stainless-steel sealed pan with a gold-plated copper seal (which can withstand 150 atm of pressure before rupturing and has a capacity of 30 mL). These measurements were carried out in a Pyris 1 calorimeter (PerkineElmer) under a temperature scan rate of 1 C min1. The weight was constant in all cases, indicating that no leaks occurred during the experiments. 3. Result and discussion As-synthesized Li[Ni0.7Mn0.3]O2 is crystallized into typical aNaFeO2 structure with R3 m space group (Fig. 1a). Surface modified SeLi[Ni0.7Mn0.3]O2 and SP-Li[Ni0.7Mn0.3]O2 also show the same crystal structure as that of the bare material (Fig. 1aec). Rietveld refinement of XRD data indicated that the presence of Ni2þ in Li
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Fig. 1. Rietveld refinement result of XRD data for (a) bare, (b) silica-coated and (c) silicon phosphate-coated Li[Ni0.7Mn0.3]O2; electrochemically delithiated (d) bare, (e) silica-coated and (f) silicon phosphate-coated Li0.3[Ni0.7Mn0.3]O2.
layer is approximately 3.9% and that the cation mixing does not vary even after the coating with silica and silicon phosphate. Also, variation in the calculated lattice parameters is negligible after the coating process. Electrochemical delithiation was carried out using these samples to avoid change derived during chemical delithiation, often resulting in morphological degradation [35]. The electrodes are first charged to 4.3 V, delivering a charge capacity of approximately 195 mAh g1, which can be expressed as Li0.3[Ni0.7Mn0.3]O2 (Fig. 2). The delivered discharge capacity reaches 170 mAh g1, which is much improved than that of our prior report, presumably due to minimization of cation mixing in the crystal structure [33]. The charged electrodes were rinsed with salt-free dimethyl carbonate for a week, in an Ar filled glove box. And, the XRD patterns of delithiated sample were calibrated using the internal standard of the Al current collector (Fig. 1def). Similar to the fresh materials, Rietveld refinement of XRD data for these delithiated electrodes are carried out by assuming R3 m space group. The original hexagonal structure is kept after the electrochemical delithiation for all samples. Refined data indicate contraction of the a-axis due to oxidation of transition metals while expansion of the c-axis resulting from increased ionicity in the oxide lattice (Table 1). Occupation of Ni2þ in Li layer also slightly increased in comparison with those of fresh materials, presumably due to migration of Ni
element on the electrochemical delithiation. Fig. 3 shows the surface morphology of the fresh and electrochemically delithiated bare, S- and SP-coated Li[Ni0.7Mn0.3]O2
Fig. 2. First charge curves of bare, silica-coated, silicon phosphate-coated Li [Ni0.7Mn0.3]O2.
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Table 1 Rietveld refinement result of XRD data for bare, silicon-coated and silicon phosphate-coated Lix[Ni0.7Mn0.3]O2 (x ¼ 1, 0.3).
Bare Li[Ni0.7Mn0.3]O2 Silica coated-Li[Ni0.7Mn0.3]O2 Silicon phosphate coated-Li[Ni0.7Mn0.3]O2 Bare Li0.3[Ni0.7Mn0.3]O2 Silica coated-Li0.3[Ni0.7Mn0.3]O2 Silicon phosphate coated-Li0.3[Ni0.7Mn0.3]O2
a-axis/Å
c-axis/Å
Ni2þ in Li layer/%
Rwp/%
2.8831(1) 2.8830(1) 2.8833(1)
14.2568(4) 14.2569(4) 14.2566(3)
3.9(3) 3.9(3) 3.9(2)
12.5 12.4 14.0
2.8381(2) 2.8385(1) 2.8384(1)
14.4035(4) 14.4033(4) 14.4031(2)
5.8(2) 5.8(3) 5.8(3)
10.7 12.9 12.7
particles. The fresh bare particles have smooth surfaces, and there are no additional layers on the surface (Fig. 3a). For S- and SPcoated particles, in contrast, surface layers with thickness about 10 nm are noticed on the surface of the particles (Fig. 3b and c). Even after the electrochemical delithiation up to Li0.3[Ni0.7Mn0.3]O2, the smooth surfaces are still observed for the coated particles (Fig. 3e and f), while the delithiated bare particles seems to be somehow damaged during the delithiation (Fig. 3d), which demonstrates that the presence of coating layer is effective in the protection of the outer surface during migration of Liþ ions. The surfaces for the electrochemically delithiated bare and coated electrodes are analyzed by ToF-SIMS (Fig. 4). In the highlighted ToF-SIMS spectra, a Si2Oþ (m ¼ 71.94) fragment is detected on the coated electrodes whereas the fragment is absent for the bare. Since the silicon phosphate consists of SiePeO bond, the appearance of Si2Oþ fragment is reasonable. And SiPOþ (m ¼ 74.94) fragment is found for the SP-coated electrode (bottom of Fig. 4b). These data coincide with the TEM results that the coating layer remains on the surface of active materials after the electrochemical
delithiation. TGA investigation was carried out for the electrochemically delithiated bare, S-, and SP-coated powders (Fig. 5). The charged powders show distinct differences in weight variation on heating. Since the powders were dried under vacuum at 80 C for a day, the weight loss associated with evaporation of water is negligible up to 150 C for all samples. Then, the first drastic weight loss begins from 200 C and is continued to 400 C, showing 9.7 wt. % of weight loss for the delithiated bare, which results from oxygen release from the parent material [29]. The second weight loss is observed in the temperature range of 400e550 C (13.2 wt. %), after which the weight keeps constant up to 600 C. For SeLi0.3[Ni0.7Mn0.3]O2, the initiation temperature of the drastic weight loss is somehow delayed to 240 C, after which the weight loss is progressed to 460 C with 8.8 wt. % of oxygen loss. The weight further decreases to 600 C approximately 8.5 wt. %. A similar tendency is observed for the SP-Li0.3[Ni0.7Mn0.3]O2. Namely, the initiation temperature of the drastic weight loss is further delayed to 260 C, and the weight loss is progressed relatively slow to 460 C (5.3 wt. %) and then accelerated to 510 C (3.3%), which belongs to the first stage of the weight loss process. In comparison with the weight loss process for S- and SP-Li0.3[Ni0.7Mn0.3]O2, it is obvious that the silicon phosphate coating (5.5 wt. %) seems to be more effective in suppressing weight loss than silica coating to 460 C (8.8 wt. %) because of less evolution of oxygen from the crystal structure of active materials. Further weight loss approximates to 7 wt. % in the temperature range of 510e600 C, which is similar to the silica coating. As a result, it is clear that oxygen release is more dominant and starts relatively at lows temperature for the bare materials, while the presence of a coating layer obviously delays the oxygen release process upon heating. To further correlate the weight loss to the structural change from the charged materials upon heating, in-situ HT-XRD
Fig. 3. TEM images for (a) bare, (b) silica-coated and (c) silicon phosphate-coated Li[Ni0.7Mn0.3]O2; electrochemically delithiated (d) bare, (e) silica-coated and (f) silicon phosphatecoated Li0.3[Ni0.7Mn0.3]O2.
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Fig. 4. ToF-SIMS results of bare, silica-coated and silicon phosphate-coated Lix[Ni0.7Mn0.3]O2 (x ¼ 1 (a, b); x ¼ 0.3 (c, d)). (top: bare, middle: silica-coated material, bottom: silicon phosphate-coated material).
measurements were carried out for the electrochemically delithiated bare, S-, and SP- Li0.3[Ni0.7Mn0.3]O2 (Fig. 6aec). The series of patterns were collected at an interval of 10 C in the temperature range of 150e300 C. Otherwise, the interval was 50 C. Bare, S- and SP-Li0.3[Ni0.7Mn0.3]O2 have a stable rhombohedral structure at room temperature. Although the majority of crystal structure remains as the rhombohedral structure, the resulting peak positions are shifted and the intensity increases or decreases to certain temperatures; namely, for bare to 200 C, for SeLi0.3Ni0.7Mn0.3O2 to 240 C, and SP-Li0.3Ni0.7Mn0.3O2 to 260 C. Also, the relative intensity ratio of I(003)/I(104) shows that the occupation of transition metal ions in the Li layers decreases gradually without formation of new phases in the same temperature range: bare: 1.03, SeLi0.3Ni0.7Mn0.3O2: 1.22, and SP-Li0.3Ni0.7Mn0.3O2: 1.29 at 50 C; bare: 0.85, SeLi0.3Ni0.7Mn0.3O2: 0.94, SP-Li0.3Ni0.7Mn0.3O2: 0.95 at 150 C; and bare: 0.68, SeLi0.3Ni0.7Mn0.3O2: 0.79, and SPLi0.3Ni0.7Mn0.3O2: 0.88 at 200 C. Lowering of the ratio is obvious with increasing the heating temperature, meaning that migration of transition metal, particularly nickel, to the Li layers is accelerated from the octahedral 3b sites to vacant 3a sites via tetrahedral 6c sites in the Li layer, which is face-shared with the 3a and 3b sites,
Fig. 5. TGA curves of bare, silica-coated and silicon phosphate-coated Li0.3[Ni0.7Mn0.3] O2.
with increasing temperature. In this charged state, in which both Ni and Mn are oxidized to tetravalent, it is likely that energetically unstable Ni4þ (0.48 Å) relative to Mn4þ, which has smaller ionic size than that of Mn4þ (0.53 Å), would readily migrate to Li layers. This behavior results in the increment of the a-axis and simultaneously affects the shrinkage of LiO2 slabs accompanying decrease in the caxis. Hence, the gradual decrease in the c/a ratio is observed to 200 C (Fig. 7). This result demonstrates that cation migration is continuously progressed in the rhombohedral symmetry without structural changes to 200 C for the bare, 240 C for the SeLi0.3Ni0.7Mn0.3O2, and 260 C for the SP-Li0.3Ni0.7Mn0.3O2. The advent of the cubic spinel structure is initiated above these temperatures. These onset temperatures are almost consistent with the starting point of weight loss in the TGA result (Fig. 5). In particular, the (108) and (110) peaks indicating a typical O3 type layer structure are merged to the (440) peak of cubic spinel structure with Fd3m space group. Since oxygen evaporation is aggressively progressed in the temperature region of 200e300 C for the bare, the new spinel (311), (400), and (440) peaks evolve fast in this temperature. And the structural transformation to the cubic spinel phase seems to be terminated at 400 C (Fig. 6a), which coincides with the first weight loss region in TGA (Fig. 5). Interestingly, these peaks shift toward lower angle, indicating thermal expansion of the crystal lattice. Above the temperature of 400 C, formation of rock salt type phase is noticed and the two phases, MO (Fm3m) and M3O4 (M: Ni and Mn) coexist up to 600 C. This temperature range belongs to the second oxygen evolution region, approximately 13.2 wt. %. During the oxygen evaporation, structural reorganization occurs via the continuous migration of cations to a more stable site such as the vacant lithium site and tetrahedral site, resulting in the formation of a simple cubic MO compound. The two phases are still observed even after cooling to room temperature. Different from the bare Li0.3Ni0.7Mn0.3O2, the SeLi0.3Ni0.7Mn0.3O2 undergoes a slow phase evolution toward cubic spinel structure in the first weight loss region, 240e460 C (Fig. 6b). At 250 C, the rhombohedral structure begins to be transformed to the cubic spinel structure, and these two phases coexist to 400 C. There are apparently no shifts of diffraction peaks to this temperature, probably due to the rearrangement of crystal structure to cubic spinel structure. As soon as the single cubic spinel phase is formed at 450 C, the diffraction peak positions move toward the
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Fig. 6. In-situ HT-XRD patterns of (a) bare, (b) silica-coated and (c) silicon phosphate-coated Li0.3[Ni0.7Mn0.3]O2 in the temperature range of 25 Ce600 C.
lower angle owing to the thermal expansion of the structure. The temperature is close to the ending point of the first weight loss in TGA (Fig. 5). The released oxygen content is about 8.8 wt. % to 460 C for the SeLi0.3Ni0.7Mn0.3O2, while the content is approximately 9.7 wt. % to 400 C for the bare Li0.3Ni0.7Mn0.3O2. Above 450 C, the advent of the simple cubic MO is observed (Fm3m) and M3O4 (M: Ni and Mn) coexist up to 600 C. In the temperature range of 450e600 C, the evolved amount of oxygen from the parent material is about 8.5 wt. %, which is lower than that of bare. Therefore, the phase transformation from the cubic spinel to the simple cubic phase is delayed. The two phases are still observed even after cooling to room temperature. These data suggest that the presence of silica coating layers suppresses oxygen evolution from the crystal structure, and this, in turn, leads the phase transition tedious. For the SP-Li0.3Ni0.7Mn0.3O2, the phase transformation toward the cubic spinel phase is observed at a slightly higher temperature (270 C, Fig. 6c) than that of the SeLi0.3Ni0.7Mn0.3O2 (250 C, Fig. 6b). The cubic spinel structure is still perceived up to 450 C due to the less amount of oxygen evolution to 510 C (5.5 wt. %), which
confirms the effect of coating that is able to delay the phase transformation triggered by the oxygen release from the parent oxide. Above 500 C, the resulting phase is a mixture of the cubic spinel and simple cubic phases. It is likely that the silicon phosphate layer is favored rather than silica layer because the series of phase transformation is delayed to higher temperature. Therefore, the phase transformation is initiated by the gradual oxygen release, which simultaneously causes rearrangement of structure via migration of metal elements, from the structure in the following sequence; (i) rhombohedral phase, (ii) coexistence rhombohedral and spinel phases, (iii) spinel phase, and (iv) coexistence of spinel and rock salt phases. As a result, the consecutive cation migration resulting from oxygen release gives rise to variation in lattice parameters (Fig. 7a). For convenience, the lattice parameters are represented as the rhombohedral symmetry even after the phase transformation to the cubic spinel structure. Drastic weight change reflects an abrupt decrease in the c/a ratio because of the advent of cubic spinel structure in the temperature of 200e450 C. Among samples, the bare material undergoes a fast structural transformation,
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Fig. 7. Variation in of c/a ratio for the electrochemically delithiated bare, silica-coated and silicon phosphate-coated Li0.3[Ni0.7Mn0.3]O2 in the temperature range of 25 Ce600 C.
terminating at about 300 C, which shows the c/a ratio close to 4.9 that is interpreted as the cubic symmetry. In comparison with the prior reports, the onset temperature of phase transformation from rhombohedral to the spinel phase for the present Li0.3[Ni0.7Mn0.3] O2 (220 C) was similar to that of Li0.3[Ni0.7Co0.15Al0.15]O2 (230 C) [18] and Li0.33[Ni0.5Co0.2Mn0.3]O2 (223 C) apart from the different ratio of cation mixing in the Li layers [35,36]. By contrast, the variation is retarded by the presence of foreign silica and silicon phosphate coating layers. Our prior report revealed the effect of AlF3 coating on Li0.35[Ni1/3Co1/3Mn1/3]O2 such that the coating layer blocks oxygen evolution during exothermic decomposition, resulting in delayed phase transition [29]. It is likely that the present silica and silicon phosphate layers play the similar role to the AlF3 layer, since the TGA and in-situ HT-XRD results are similar to the previous results even in Ni-rich Li0.3Ni0.7Mn0.3O2 composition. The electrochemically delithiated electrodes were heated at 250 C, 400 C and 600 C, whose temperatures correspond to phase transition points, in air for 1 h and quenched to room temperature for surface and ToF-SIMS analyses. Fig. 8 exhibits the surface morphology and the chemistry information for the bare, S-, and SP-Li0.3Ni0.7Mn0.3O2 powders quenched at 250 C. For the bare
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material, outer part is damaged, and it progresses toward the particle inside presumably due to the oxygen evolution from the oxide particle after the quenching at 250 C (Fig. 8a). The appearing cubic spinel phase in the in-situ XRD patterns would be ascribed to the damaged surface after the oxygen evolution. Unlike the bare particle, the S- and SP-Li0.3Ni0.7Mn0.3O2 particles present smooth and clear surfaces (Fig. 8b and d). In addition, the silica and silicon phosphate coating layers are still visible, as further evidenced from the Si2Oþ fragment and silicon phosphate layer from the SiPOþ fragment (Fig. 8c and e), indicating the presence of coating layers without thermal decomposition at this temperature. Oxygen would be aggressively removed from the oxide particle for the bare material at 250 C as evident from TGA, whereas the oxygen evolution is just initiated at the temperature for the coated samples. At 400 C (Fig. 9a), better particle stability is noticed for both coated particles (Fig. 9b and d), and the coating layers are still observed. It seems that the coated particles remained intact, although there was weight loss of approximately 6.7 wt. % for the SeLi0.3Ni0.7Mn0.3O2 and 3.7 wt. % for the SP-Li0.3Ni0.7Mn0.3O2. Less deterioration observed in those particles is reasonable because the quenching was done at the intermediate temperature of the phase transformation. Additional information is given by the ToF-SIMS spectra that the silica, as the SiOþ fragment, and silicon phosphate, as the SiPOþ fragment, coating layers are not decomposed but still appear on the outermost surface (Fig. 9c and e). At 600 C, the crystal structures of the three delithiated materials are stabilized into the simple cubic structure as a result of the two-step oxygen evolution from the parent particles. Unlike the bare material that starts to form the rock salt phase from 400 C, the advent of the rock salt phase is delayed for the coated materials: 450 C for the SeLi0.3Ni0.7Mn0.3O2 and 500 C for the SPLi0.3Ni0.7Mn0.3O2, which would result from the less damage of the delithiated active materials even at 600 C. In addition, coating media remained on the outermost surface of the active material as the presence of coating layers are clearly observed on the TEM images and ToF-SIMS spectra (Fig. 10bee). Note that the coating layers are very thin, below 10 nm (Fig. 3). The coating layers probably block the oxygen evolution from the oxide lattice. As evident from the TGA, XRD, TEM, and ToF-SIMS results, the SP-Li0.3Ni0.7Mn0.3O2 experiences slow phase
Fig. 8. TEM images and ToF-SIMS results of the bare (a), silica coated (b, c) and silicon phosphate coated Li0.3[Ni0.7Mn0.3]O2 (d, e) quenched at 250 C. (in ToF-SIMS result, upper: Si2Oþ fragment, bottom: SiPOþ fragment).
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Fig. 9. TEM images and ToF-SIMS result of the bare (a), silica coated (b, c) and silicon phosphate coated Li0.3[Ni0.7Mn0.3]O2 (d, e) quenched at 400 C. (in ToF-SIMS result, upper: Si2Oþ fragment, bottom: SiPOþ fragment).
transformation due to less evaporation of oxygen from the parent oxide. It is thought that the presence of PeO covalent bond may assist to form the coating layer much solid relative to the silica coating, and this, in turn, would further delay the oxygen evaporation from the delithiated Li0.3Ni0.7Mn0.3O2. Fig. 11 represents DSC data for the electrochemically delithiated bare, S-, and SP-Li0.3Ni0.7Mn0.3O2 in the presence of electrolyte. The onset temperatures of exothermic reaction coincide with the TGA and in-situ HT-XRD data for the three samples, in which the cubic spinel phase begins to be produced as a result of oxygen loss from the rhombohedral oxide lattice. The main exothermic decomposition occurs at a somehow elevated temperature by the silica and silicon phosphate coating, and the generated heat is also dependent on the coating layer. Specifically, less evolution of oxygen brings about less exothermic heat as follows: bare (~614.4 J g1), silica coated-material (~498.6 J g1), and silicon phosphate-coated
material (~432.7 J g1). It is most likely that the coating layers block the oxygen release from the active material, and it, in turn, affects the shift of the main exothermic reaction to a higher temperature. After the reaction, the majority of the resulting phase is composed of the cubic spinel phase as confirmed in the in-situ HTXRD. Therefore, the presence of coating layers on the surface of the delithiated active materials is responsible for the improved thermal properties. 4. Conclusion The thermal behavior of the delithiated bare, silica-coated and silicon phosphate-coated Li0.3[Ni0.7Mn0.3]O2 is investigated. Oxygen release from the parent oxide materials induces phase transformation in a sequence of rhombohedral to cubic spinel and finally the simple cubic phase. Applying the coating layers on the
Fig. 10. TEM images and ToF-SIMS result of the bare (a), silica coated (b, c) and silicon phosphate coated Li0.3[Ni0.7Mn0.3]O2 (d, e) quenched at 600 C. (in ToF-SIMS result, upper: Si2Oþ fragment, bottom: SiPOþ fragment).
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References
Fig. 11. Comparison of DSC traces of the bare, silica coated and silicon phosphate coated Li0.3[Ni0.7Mn0.3]O2.
surface of cathode materials, the onset temperature of the exothermic decomposition is retard over 40 C compared to the bare material because the oxygen loss is suppressed by the coating layers. In particular, silicon phosphate coating is effective in delaying the exothermic reaction relative to the oxide coating due to the sturdiness of the coating layer supported by the presence of covalent bond of silicon phosphate. Therefore, the role of the coating layer and the related mechanism behind the thermal properties are comprehensively understood by means of TGA, TEM, in-situ HT-XRD, and ToF-SIMS, which are used to interpret the DSC behavior. Acknowledgments The authors thank Miwa Watanabe, Iwate University, for her helpful assistance in the experimental work. This research was supported by Basic Science Research Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Education, Science and Technology (NRF-2009-0093467). This work was also supported by MKE/KEIT (10041856 and 10041094). This research was supported by Basic Science Research Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Education, Science and Technology (2014R1A2A1A11051197)
[1] B. Scrosati, J. Garche, J. Power Sources 195 (2010) 2419. [2] M. Winter, J.O. Besenhard, M.E. Spahr, P. Novak, Adv. Mater. 10 (1998) 725. [3] V. Etacheri, R. Marom, R. Elazari, G. Salitra, D. Aurbach, Energy Environ. Sci. 4 (2011) 3243. [4] L. Lu, X. Han, J. Li, J. Hua, M. Ouyang, J. Power Sources 226 (2013) 272. [5] M.M. Thackeray, C. Wolverton, E.D. Issacs, Energy Environ. Sci. 5 (2015) 7854. [6] Y.-K. Sun, Z. Chen, H.-J. Noh, D.-J. Lee, H.-G. Jung, Y. Ren, S. Wang, C.-S. Yoon, S.-T. Myung, K. Amine, Nat. Mater. 11 (2012) 942. [7] P.Y. Liao, J.G. Duh, S.R. Sheen, J. Electrochem. Soc. 152 (2005) A1695. [8] A. Wuersig, W. Scheifele, P. Nov ak, J. Electrochem. Soc. 154 (2007) A449. [9] X. Feng, J. Sun, M. Ouyang, X. He, L. Lu, X. Han, M. Fang, H. Peng, J. Power Sources 272 (2014) 457. [10] J. Lamb, C.J. Orendorff, K. Amine, G. Krumdick, Z. Zhang, L. Zhang, A.S. Gozdz, J. Power Sources 247 (2014) 1011. [11] Z. Chen, Y. Ren, E.J. Lee, C. Johnson, Y. Qin, K. Amine, Adv. Energy Mater. 6 (2013) 729. [12] L. Wu, K.W. Nam, X. Wang, Y. Zhou, J.C. Zheng, X.Q. Yang, Y. Zhu, Chem. Mater. 23 (2011) 3953. [13] K.W. Nam, S.M. Bak, E. Hu, Z. Yu, Y. Zhou, X. Wang, L. Wu, Y. Zhu, K.Y. Chung, X.Q. Yang, Adv. Funct. Mater. 23 (2013) 1047. [14] J. Molenda, P. Wilk, J. Marzec, Solid State Ionics 146 (2002) 73. [15] B.H. Kim, J.H. Kim, I.H. Kwon, M.Y. Song, Ceram. Int. 33 (2007) 837. [16] Z. Liu, H. Zhen, Y. Kim, C. Liang, J. Power Sources 196 (2011) 10201. [17] N. Yabuuchi, Y.T. Kim, H.H. Li, S.H. Yang, Chem. Mater. 20 (2008) 4936. [18] M. Guilmard, A. Rougier, A. Grune, L. Croguennec, C. Delmas, J. Power Sources 115 (2003) 305. [19] A. Rougier, I. Saadoune, P. Gravereau, P. Willmann, C. Delmas, Solid State Ionics 90 (1996) 83. [20] S.T. Myung, S. Komaba, K. Kurihara, K. Hosoya, N. Kumagai, Y.K. Sun, I. Nakai, M. Yonemura, T. Kamiyama, Chem. Mater. 18 (2006) 1658. [21] J. Jiang, K.W. Eberman, L.J. Krause, J.R. Dahn, J. Electrochem. Soc. 152 (2005) A566. [22] S.T. Myung, K. Izumi, S. Komaba, Y.K. Sun, H. Yashiro, N. Kumagai, Chem. Mater. 17 (2005) 2427. [23] S.T. Myung, K. Izumi, S. Komaba, H. Yashiro, H.J. Bang, Y.K. Sun, N. Kumagai, J. Phys. Chem. C 111 (2007) 4061. [24] M. Bettge, Y. Li, B. Sankaran, N.D. Rago, T. Spila, R.T. Haasch, I. Petrov, D.P. Abraham, J. Power Sources 233 (2013) 346. [25] K.S. Lee, S.T. Myung, K. Amine, H. Yashiro, Y.K. Sun, J. Mater. Chem. 19 (2009) 1995. [26] K.J. Rosina, M. Jiang, D.L. Zeng, E. Salager, A.S. Best, C.P. Grey, J. Mater. Chem. 22 (2012) 20602. [27] A.T. Appapillai, A.N. Mansour, J.-P. Cho, Y. Shao-Horn, Chem. Mater. 19 (2007) 5748. [28] D.J. Lee, B. Scrosati, Y.K. Sun, J. Power Sources 196 (2011) 7745. [29] S.T. Myung, K.S. Lee, C.S. Yoon, Y.K. Sun, K. Amine, H. Yashiro, J. Phys. Chem. C 114 (2010) 4710. [30] K.H. Teoh, S. Ramesh, A.K. Arof, J. Solid State Electrochem. 16 (2012) 3165. [31] D. Praveen, S.V. Bhat, R. Damle, Ionics 19 (2013) 1375. [32] K.S. Lee, S.T. Myung, J.S. Moon, Y.K. Sun, J. Electrochem. Soc. 161 (2014) A920. [33] D.H. Cho, C.H. Jo, W.S. Cho, Y.J. Kim, H. Yashiro, Y.K. Sun, S.T. Myung, J. Electrochem. Soc. 161 (2014) A920. [34] T. Roisnel, J. Rodriguez-Carjaval, Fullprof Manual, Institut Laue-Langevin, Grenoble, France, 2002. [35] Y.H. Cho, D.H. Jang, J.B. Yoon, H.C. Kim, T.K. Ahn, K.W. Nam, Y.E. Sung, W.S. Kim, Y.S. Lee, X.Q. Yang, W.S. Yoon, J. Alloys Compd. 562 (2013) 219. [36] M. Guilmard, L. Croguennec, C. Delmas, Chem. Mater. 15 (2003) 4484.