Diamond-shaped Fe₂O₃@C₁₈H₃₄O₂ core–shell nanostructures as anodes for lithium ion batteries with high over capacity

Abstract

The over capacity of lithium ion batteries with metal-oxides anode materials is usually regarded as originating from the low-voltage decomposition of the electrolyte and subsequent formation of a gel-like polymer layer deposited on the metal-oxides surfaces. In this work, we report a high over capacity value of 1800 mA h g⁻¹ after 350^th charge–discharge cycles for Fe₂O₃-made lithium ion batteries. It is found that the capacitive nature of the designed Fe₂O₃@C₁₈H₃₄O₂ nanostructure not only contributes to the large observed excess in capacity, but also results in unique rate capabilities. Thus, a capacitive model is proposed to outline a plausible mechanism to explain these electrochemical findings, and it is anticipated that this paper will shed some new light on future design of the next generation of lithium ion batteries.

---

1. Introduction

Energy storage will be more important in the future than at any time in the past. Among the myriad of energy-storage technologies, lithium-ion battery technology is undergoing rapid expansion, now representing the largest segment of the portable battery industry and dominating the computer, cell phone and camera power source industry.^1,2 However, the present lithium-ion battery industry is confronted with formidable challenges, such as high cost, unsafety, environmental contamination, limited scale, and low power in terms of specific energy (energy per unit weight) and energy density (energy per unit volume).^1,3–5

In order to address these challenges to the greatest extent, extensive studies have been devoted to iron oxides, especially trivalent iron oxides Fe₂O₃.^6,7 The conversion reactions of these interstitial-free 3d metal oxides with structures unsuitable for intercalation chemistry have been shown to exhibit large, rechargeable specific capacities (as high as 1005 mA h g⁻¹) in cells with lithium, and hence being regarded as one of the most potential candidates for anode materials. Unfortunately, although the reversible specific energy can be greatly enhanced in these materials, the overall cell energy density is enhanced substantially less than what would be expected by specific energy alone. That being the case, reducing the particle size of Fe₂O₃ active materials, as a popular way to improve their rate capabilities due to small diffusion lengths, would result in a low tap density which will further reduce their energy density. One direct and effectual avenue to overcome this dilemma is to greatly enhance their specific capacity, even surmounting the theoretical capacity (namely over capacity), with the volume of electrode materials held nearly constant. Despite many phenomenological results of over capacity reported in the literature,^8–10 the rational design and controllable synthesis of Fe₂O₃ nanomaterials with large excesses in capacity remain as significant challenges, and are still in its very early stage. It is hence highly desirable to unlock the tremendous potential of over capacity currently at its discovery phase.

In this work, we fabricate a Fe₂O₃@C₁₈H₃₄O₂ core–shell nanostructure, in which the C₁₈H₃₄O₂ shell is designed to form double charge layer with the Fe₂O₃ core, and hence inducing striking capacitive behaviors to enhance the specific capacity of Fe₂O₃ anode. The observed over capacity value of 1800 mA h g⁻¹ after 350^th cycles is, to our knowledge, the highest among that of any metal oxides. More importantly, such design also endows the Fe₂O₃ anode with unique rate capabilities, which enable us to gain deep insight into its electrochemical performances.

2. Experimental

2.1. Preparation of Fe₂O₃@C₁₈H₃₄O₂ core–shell nanostructure

The Fe₂O₃@C₁₈H₃₄O₂ core–shell nanostructure was fabricated according to our previous work¹¹ with little modification. In a typical procedure, 1.92 g of sodium oleate (6 mmol), 20 mL of oleic acid (C₁₈H₃₄O₂) and 40 mL of ethanol were mixed by vigorously stirring, to which 10 mL of Fe(NO₃)₃·9H₂O aqueous solution (0.2 mol L⁻¹) was added at room temperature. The dark red solution mixture was then transferred into a Teflon-lined stainless-steel autoclave with a capacity of 100 mL for hydrothermal treatment at 180 °C for 20 hours. The as-obtained precipitate was repeatedly washed with n-hexane, ethanol and deionized water, and finally dried at 60 °C for 4 hours.

2.2. Characterization

The XRD patterns were recorded on a powder X-ray diffractometer (Rigaku D/max-rA) equipped with a rotating anode and a Cu Kα₁ radiation source (λ = 1.5406 Å) at a step width of 0.02°. Transmission electron microscopy (TEM) was performed on the JEOL 2010 TEM with operating voltage at 200 kV. The high-resolution transmission electron microscopy (HRTEM) was conducted using a field emission gun (FEG) JEOL 2010F microscope with a point resolution of 0.19 nm. Fourier transform infrared (FTIR) spectrometry was performed on KBr disks of the powdered samples using a BRUKER TENSOR 27 FTIR spectrometer.

2.3. Electrochemical measurements

The electrochemical measurements were carried out using 2032 coin-type cells with pure lithium metal as the counter and reference electrodes at two different temperatures of 25 °C and 55 °C. The working electrode consists of active material (as-synthesized Fe₂O₃@C₁₈H₃₄O₂ powder), conductive agent (carbon black), and sodium carboxymethyl cellulose (CMC, 800–1200 mPa s) binder in a weight ratio of 80 : 10 : 10, using deionized water as the dispersion medium. The weights of Fe₂O₃@C₁₈H₃₄O₂ powder loaded onto one electrode are 0.48, 1.80 and 3.78 mg. The mixture was spread on a Cu foil and dried under vacuum at 120 °C for 8 hours. The electrolyte used was 1.0 mol L⁻¹ LiPF₆ in a mixture of ethylene carbonate (EC) and dimethyl carbonate (DMC) (1 : 1 by volume). Cell assembly was carried out in an Ar-filled glove box. The cells were cycled at different current rates of 0.2, 0.5, 1, 2 and 5 C between 0.01 and 3 V using a LAND battery tester. Electrochemical impedance spectra (EIS) measurements were carried out at room temperature using CHI660D electrochemical workstation over a wide frequency range from 100 kHz to 10 mHz with an ac perturbation voltage of 10 mV. The impedance data were fitted using the Autolab Nova 1.8 software. The cyclic voltammetry (CV) measurements were performed at different voltage ranges of 0.02–2 V and 0.02–1.5 V at different scan rates of 0.1, 0.5, 1, 1.5, 2, 2.5, 3, 3.5, 4, 4.5, 5, 5.5 and 6 mV s⁻¹. The specific capacity was approximately calculated based on the mass of Fe₂O₃ alone.

3. Results and discussion

3.1. Morphology and microstructure of Fe₂O₃@C₁₈H₃₄O₂

The chemical composition of the as-prepared product is shown in Fig. 1a, in which all the diffraction peaks can be unambiguously indexed to the hexagonal structure of hematite (JCPDS no. 33-0664) and the cubic structure of maghemite (JCPDS no. 39-1346). A broad and weak apophysis at low diffraction angles indicates the remnant of amorphous C₁₈H₃₄O₂ phase, even after extensive washing, which can be further confirmed by the FTIR spectrum in Fig. 1b. The two bands of 2923 and 2852 cm⁻¹ represent the aliphatic alkyl groups of C₁₈H₃₄O₂, revealing the chemisorption of oleic group on the surface of Fe₂O₃ particles.¹² And, the bands at 1625 and 1431 cm⁻¹ can be assigned to the vibration of carboxylate groups and C–H bond, respectively. The peaks at 555 and 480 cm⁻¹ are essential features for Fe–O bond in Fe₂O₃.

Fig. 1 (a) XRD pattern and (b) FTIR spectrum of the as-obtained product.

The TEM image in Fig. 2a shows the product is of diamond shape, with average side length of 15 nm. As is evident from the periodic lattice fringes across all the selected diamond-shaped particles (Fig. 2b–h), each of these particles is a monocrystal regardless of their phase composition. More importantly, a surface C₁₈H₃₄O₂ shell, being of amorphous characteristic with different thickness, was also clearly observed. The dissimilar electronic density of Fe₂O₃ and C₁₈H₃₄O₂ phases allow clear differentiation of core–shell structure of the nanoparticles. Such observation further confirms the XRD and FTIR analyses in Fig. 1. Although the surface coated C₁₈H₃₄O₂ layer is always avoidless and normally undesired in most C₁₈H₃₄O₂-participated chemical syntheses, such an experimental drawback can thereafter turn into an advantage to generate an uniform double charge layer upon charge–discharge cycling as will be discussed in the forthcoming section.

Fig. 2 (a) TEM and (b–h) HRTEM images of the as-obtained product. The lattice spacings of α-Fe₂O₃ and γ-Fe₂O₃ components are marked with white and black, respectively.

3.2. Electrochemical performances

The charge–discharge cycling was carried out for Fe₂O₃@C₁₈H₃₄O₂ (0.48 mg)/Li cell in the voltage window of 0.01–3 V (vs. Li/Li⁺) at a current density of 200 mA g⁻¹ (0.2 C). The voltage versus capacity profiles are shown in Fig. 3a. During the 1^st cycle, the first plateau appears at 1.6 V, followed by a smooth voltage drop to 1.2 V, wherein a canted plateau occurs and hence follows by another voltage drop to ∼0.9 V. This has been reported as a feature for lithium insertion in Fe₂O₃ to form Li_xFe₂O₃.^8,13 The two different plateau voltages illustrate the different lithium insertion potentials as for α-Fe₂O₃ and γ-Fe₂O₃ phases in the electrode. Then, a third flat voltage plateau at about 0.9 V, probably ascribed to the formation of solid electrolyte interphase (SEI) layer and the reduction of Fe³⁺ to Fe⁰,^8,13 sets in and continues until a capacity of ∼1700 mA h g⁻¹ is reached, followed by a gradual drop in voltage until the end of discharge. The capacity retention during the initial 30 charge–discharge cycles is rather disappointing (Fig. 3b), in that the specific capacity value sharply drops down from 2840 to 660 mA h g⁻¹. Upon further cycling, the capacity retention curve levels off between 30 and 60 cycles, and thereafter the specific capacity value ever-increases to 1800 mA h g⁻¹ after 350^th charge–discharge cycles (Fig. 3b).

Fig. 3 (a) Charge–discharge voltage profiles of Fe₂O₃@C₁₈H₃₄O₂ (0.48 mg)/Li cell cycled between 0.01 and 3 V at 0.2 C rate. (b) 25 °C capacity retention of Fe₂O₃@C₁₈H₃₄O₂ (0.48 mg)/Li (solid square) and Fe₂O₃ (1.56 mg)/Li (solid circle) cells cycled between 0.01 and 3 V at 0.2 C rate. (c) Schematic illustration of a possible capacitive model. (d) Weight and temperature effects of Fe₂O₃@C₁₈H₃₄O₂/Li cells on the capacity retention cycled between 0.01 and 3 V at 0.2 C rate.

Since Fe₂O₃ neither exhibits interstitial sites for lithium, nor alloys with Li, a two-step mechanism was popular to explain the observed high specific energy.⁹ The first and best understood mechanism, namely conversion reaction, entails the decomposition of Fe₂O₃ into Fe nanoparticles embedded into a Li₂O matrix, which converts back to Fe₂O₃ nanoparticles upon subsequent oxidation, leading to the overall reaction Fe₂O₃ + 6Li ↔ 3Li₂O + 2Fe.¹⁴ Such a conversion reaction was demonstrated to be thermodynamically feasible and kinetically favored by the presence of Fe nanoparticles formed during the first reduction step. The second one was supposed to be associated to the growth and dissolution of a gel-like polymer layer around Fe₂O₃ particles upon subsequent cycles. A charge transfer process of electrons to the alkyl carbonate molecules forms radical anions, which are stabilized by Li-ions that are adsorbed on the Fe particles. Upon reduction, a piling of these stabilized radical ions, which should give back solvent or by-product molecules at higher potentials upon oxidation, thus give the reversible capacity. It appears that the reduction of Fe₂O₃ by lithium consists, during the first discharge, of these two separate and distinct processes that become highly intermixed. Upon long cyclings, the Fe₂O₃ ↔ Fe process, initially corresponding to the main part of the overall cell capacity, gradually vanishes with a predominance of the polymer-like process over the conversion process. As such, the capacity retention curve in Fig. 3b progressively lowers down and levels off around 30–60 cycles, indicating that the momentum of capacity loss minted by the disappearing conversion process is effectively contained by the progressively enhancing polymer-like process.

Upon further cyclings, the capacity retention in Fig. 3b unceasingly increases, even approaches 1800 mA h g⁻¹ beyond 350^th cycles, which is well above the theoretical capacity of 1005 mA h g⁻¹ for Fe₂O₃ anode materials. The over capacity has been previously attributed to the growth of the aforementioned electrochemical gel-like polymer layer.^8–10 However, our observed capacity value of 1800 mA h g⁻¹ greatly surmounts any other reported capacity values in polymer-assisted metal-oxide electrode materials. That being the case, other mechanisms relevant to over capacity should be considered. Firstly, the intentionally designed Fe₂O₃@C₁₈H₃₄O₂ core–shell nanostructure bears striking similarities with electrical double-layer capacitors, in which the electrons are blocked out, and thus accumulating on the superficial layer of C₁₈H₃₄O₂ shell (Fig. 3c). Secondly, the adsorptive mechanism (reversible interfacial reaction), the capacity of which depends on the grain size in the first place and indeed resorts on the presence of nanoparticles, probably has critical impact on the observed over capacity. Different possibilities of interfacial reactions have been mentioned in the literature,¹⁵ among which the charge separation at phase boundaries (i.e., interfacial charging), akin to the characteristic of a pseudocapacitors, has been suggested to be very reversible unlike Li storage in the SEI layer that might also occur. Such interfacial reaction associated with charge separation is also depicted in Fig. 3c, wherein the electrons can be injected through the outer double charge layer into the Fe₂O₃ electrode matrix under a certain current rate.

Similar variation trend of capacity retention can be found in Fig. 3d. The characteristic of large decline and then rise of the capacity retention curve is a common feature as for Fe₂O₃@C₁₈H₃₄O₂ (1.80 mg)/Li and Fe₂O₃@C₁₈H₃₄O₂ (3.78 mg)/Li cells at 25 °C and 55 °C. Also, such capacity retention differs significantly from that of uncoated 25 nm sized Fe₂O₃ particles (Fig. 3b), and hence revealing the important role of C₁₈H₃₄O₂ shell.

EIS measurements of the fresh and 353^rd-cycled cells were carried out to evidence the presence of surface double charge layer and interfacial charging event. In the case of the Fe₂O₃ fresh cell (Fig. 4a), the sole semicircle at high frequency is related to active electron transfer resistance on the electrode interface, whereas the 2^nd semicircle at lower frequency unique for the 353^rd-cycled cell (Fig. 4b) is related to the Li absorption/desorption process, that is, charge transfer at interface. The appearance of the 2^nd semicircle after long cyclings reveals the interface between passivation films (including SEI layer and C₁₈H₃₄O₂ shell) and active material, which is activated by formation or decomposition of Fe₂O₃ or the mixture of Li₂O and Fe at the initial cycles, is becoming progressively accessible for Li⁺ ion transfer as increasing the cycle number.

Fig. 4 Family of Nyquist plots measured for (a) fresh cell and (b) 353^rd-cycled cell. Insets: high frequency region of corresponding Nyquist plots.

In order to provide direct and solid evidence for the presence of capacitive behaviors induced by surface double charge layer and interfacial charging event, a set of CV measurements were performed. Given that the main redox peaks appear around 1.8 V (inset of Fig. 5a), a series of CV curves regarding the 353^rd-cycled Fe₂O₃@C₁₈H₃₄O₂ (0.48 mg)/Li cell were measured by lowering the discharge cutoff voltage from 2.0 V to 1.5 V (Fig. 5a). In this case, the main redox peaks are no longer present; surprisingly, the anodic and cathodic polarization curves do not superimpose, but exhibit a quasi-rectangularly shaped profile. Such a mirror-image relation between anodic and cathodic current responses profiles in CV curves is usually reminiscent of a capacitive behavior.¹⁶

Fig. 5 CV curves for 353^rd-cycled Fe₂O₃@C₁₈H₃₄O₂ (0.48 mg)/Li cell in voltage range of 0.02–1.5 V (a) at a scan rate of 0.5 mV s⁻¹ and (b) at scan rates varying from 0.1 to 6 mV s⁻¹. The insets show (a) CV curves measured for fresh cell in voltage range of 0.02–2.0 V at a scan rate of 0.5 mV s⁻¹ and (b) the variation of the current (taken as constant over 0.5 to 1 V voltage range at scan rates of 0.1–2.5 mV s⁻¹) as a function of the scan rate (solid squares), fitted by the square root power law and linear lines.

To further test this finding we measured the variation of current (I) as a function of the scanning voltage rate (dV/dt), as shown in the inset of Fig. 5b. The current here is taken as constant over 0.5 to 1 V voltage range at scan rates of 0.1–2.5 mV s⁻¹. As opposed to the usual square root power law characteristic of a faradic process for electrical double-layer capacitors and linear behavior for pseudocapacitors, we observed a compromise variation of I as a function of dV/dt, indicating a blend of electrical double-layer capacitive and pseudocapacitive behaviors (deriving from surface double charge layer and interfacial charging event, respectively) for the reversible low voltage process in Fe₂O₃@C₁₈H₃₄O₂ (0.48 mg)/Li cell. It is well-known that if the anodic currents are proportional to the scan rate (i.e., I ∝ dV/dt), then the redox process is believed to be surface-confined, yielding pseudocapacitance-type behavior; otherwise, if the anodic currents are proportional to the square root of the scan rates (I ∝ (dV/dt)^1/2), then the redox process is considered to be diffusion-controlled, yielding electrical double-layer capacitance-type behavior. Evidently, the I ∝ (dV/dt)^a curve in the inset of Fig. 5b, where a is a critical exponent with a value of ∼0.7 in our case, strongly confirms the coexistence of electrical double-layer capacitive and pseudocapacitive behaviors. Notably, the shape of the CV profiles in Fig. 5b changes from a quasi-rectangle at low scan rates (0.1–2.5 mV s⁻¹) into a spindle at high scan rates (3–6 mV s⁻¹), further confirming the remarkable contribution of these two kinds of capacitive behaviors as increasing the scan rates.

To understand more profoundly the electrochemical performances of Fe₂O₃@C₁₈H₃₄O₂ core–shell nanostructure, it is indispensable to study its rate capabilities. Fig. 6a shows the rate capabilities of 300^th-cycled Fe₂O₃@C₁₈H₃₄O₂ (0.48 mg)/Li cell (25 °C) at different current rates (0.2–0.5–1–2–5–0.2 C). There is always a capacity drop immediately after switching from a lower current rate to a higher current rate, which can be usually explained by the concentration polarization of Li⁺ ions in the Fe₂O₃ anode resulting from a diffusion limited process.¹⁷ Notably, a reversible capacity as high as ∼1020 mA h g⁻¹, barely more than the theoretical capacity of Fe₂O₃, can still be held at 1 C for 50 cycles. As the current rate increases from 1 to 2 C and then to 5 C, the capacity retention becomes worse, only exhibiting a capacity of ∼400 mA h g⁻¹ at 5 C for 50 cycles. Further recovering the current rate back to 0.2 C, the capacity value soars up to ∼1610 mA h g⁻¹ and rapidly rises to ∼1860 mA h g⁻¹ during the initial 10 cycles at 0.2 C rate. Afterwards, the capacity value abruptly drops down to 1005 mA h g⁻¹ and exhibits quite similar capacity retention with that as illustrated in Fig. 3b.

Fig. 6 (a) Rate capabilities for Fe₂O₃@C₁₈H₃₄O₂ (0.48 mg)/Li cell after 300^th cycles at different discharging rates: 0.2–0.5–1–2–5–0.2 C. (b) Nyquist plot measured for Fe₂O₃@C₁₈H₃₄O₂ (0.48 mg)/Li cell after 350^th-cycled rate capabilities. The high-medium-frequency region is shown as an inset.

In essence, the capacity drop as increasing the current rates is intimately associated with the non-efficient grinding event, which happens at high rates due to diffusion limitations.⁹ At the effect of the strain generated by the fast lithium insertion/deinsertion process at high rates, the Fe₂O₃@C₁₈H₃₄O₂ nanostructure is shattered, and hence losing its core–shell continuity (Fig. 7a and b). In this case, the previously well-formed surface double charge layer and interfacial charging zone are severely destroyed. During subsequent charge–discharge cycles, the system gradually develops sub-structures (i.e., nano-domains as depicted in Fig. 7c) within the original nanosized particles, with the lattice strain being accommodated by slippage at the domain wall boundaries. More importantly, the formerly fragmentized C₁₈H₃₄O₂ wall restores to its original shape with decreasing thickness due to ever-increasing surface area of these nano-domains as prolonging such grinding process. In these scenarios, the specific capacity of such an electrode is proportional to 1/L, L being the capacitor size. It is evident that the recovery of surface/interface double charge layer and interfacial charging zone would induce abrupt capacity rise (e.g., the observed high capacity of 1610 mA h g⁻¹ in Fig. 6a) and its further increasing trend, from 1610 to 1860 mA h g⁻¹, due to the ever-decreasing L value. As the way things going, the amount of C₁₈H₃₄O₂ phase is progressively unable to coat the enlarging surface area of nano-domains, and finally fails to form a continuous shell; that is, the capacitive behaviors of the system are destroyed again with irreparability. That being the case, the specific capacity would drop down to a lower value of 1005 mA h g⁻¹ (Fig. 6a). Upon further cyclings, the electrolyte decomposition process dominates the subsequent capacity retention and, to some extent, mitigates ill effects on the capacity fading induced by the discontinuous C₁₈H₃₄O₂ shell.

Fig. 7 (a) TEM and (b) HRTEM images of the Fe₂O₃@C₁₈H₃₄O₂ active material after 5 C cyclings. The fast Fourier transformation pattern is shown as an inset. (c) A schematic representation of the nano-domain structure within a Fe₂O₃@C₁₈H₃₄O₂ particle after high-rate cyclings.

In order to grasp further information on the charge-transport dynamics, an ac impedance measurement was conducted for a Fe₂O₃@C₁₈H₃₄O₂ (0.48 mg)/Li cell after 350^th-cycled rate capabilities. The Nyquist plot in Fig. 6b consists of two semicircles followed by a nearly linear-type frequency variation at low frequencies, as one would expect for a pure capacitance behavior. Nevertheless, by lowering our measurements down to very low frequency of 0.1 mHz, we failed to observe a drastic deviation from a straight line with the appearance of a third large semicircle masking the two previous ones, which is reminiscent of a pseudocapacitive behavior. It follows that the grinding event induced by the high-rate capability process would be very detrimental to the pseudocapacitive behavior derived from the adsorptive mechanism. By contrast, the electrical double-layer capacitive behavior survives such grinding process at the expense of getting the C₁₈H₃₄O₂ shell thinner and more discontinuous.

4. Conclusions

In summary, the electrochemical performances of the designed Fe₂O₃@C₁₈H₃₄O₂ nanostructure were systematically studied. The observed outstanding properties should be attributed to its distinct structure that offers the following benefits: (1) the coated C₁₈H₃₄O₂ shell is conducive to form surface double charge layer, which exhibits electrical double-layer capacitance behavior; (2) at certain current rates, the C₁₈H₃₄O₂ shell is beneficial to accessing electrons and hence inducing interfacial charging event, which is responsible for the observed pseudocapacitance behavior; (3) the C₁₈H₃₄O₂ nano-shell provides negligible volumetric contribution to the overall electrode, whereas it yields strikingly high over capacity of 1800 mA h g⁻¹ after 350^th cycles. As such, the overall cell energy density would be substantially enhanced.

Acknowledgements

The authors are grateful to the financial aid from the National Natural Science Foundation of China (NSFC no. 51072087) and Specialized Research Fund for the Doctoral Program of Higher Education (20113719110001).

Notes and references

1. N. S. Choi, Z. Chen, S. A. Freunberger, X. Ji, Y. K. Sun, K. Amine, G. Yushin, L. F. Nazar, J. Cho and P. G. Bruce, Angew. Chem., Int. Ed., 2012, 51, 9994.
2. Z. Wang, L. Zhou and X. W. Lou, Adv. Mater., 2012, 24, 1903.
3. M. Armand and J. M. Tarascon, Nature, 2008, 451, 652.
4. P. G. Bruce, S. A. Freunberger, L. J. Hardwick and J. M. Tarascon, Nat. Mater., 2012, 11, 19.
5. B. Scrosati, Nat. Nanotechnol., 2007, 2, 598.
6. M. V. Reddy, T. Yu, C. H. Sow, Z. X. Shen, C. T. Lim, G. V. S. Rao and B. V. R. Chowdari, Adv. Funct. Mater., 2007, 17, 2792.
7. B. Wang, J. S. Chen, H. B. Wu, Z. Wang and X. W. Lou, J. Am. Chem. Soc., 2011, 133, 17146.
8. D. Larcher, C. Masquelier, D. Bonnin, Y. Chabre, V. Masson, J. B. Leriche and J. M. Tarascon, J. Electrochem. Soc., 2003, 150, A133.
9. S. Grugeon, S. Laruelle, L. Dupont and J. M. Tarascon, Solid State Sci., 2003, 5, 895.
10. P. C. Wang, H. P. Ding, T. Bark and C. H. Chen, Electrochim. Acta, 2007, 52, 6650.
11. X. Zhang, J. Ma, N. Liu, G. Li and K. Chen, J. Electrochem. Soc., 2013, 160, A1163.
12. L. Wang and L. Gao, J. Phys. Chem. C, 2009, 113, 15914.
13. D. Larcher, D. Bonnin, R. Cortes, I. Rivals, L. Personnaz and J. M. Tarascon, J. Electrochem. Soc., 2003, 150, A1643.
14. S. Yuan, Z. Zhou and G. Li, CrystEngComm, 2011, 13, 4709.
15. J. Jamnik and J. Maier, Phys. Chem. Chem. Phys., 2003, 5, 5215.
16. S. Laruelle, S. Grugeon, P. Poizot, M. Dollé, L. Dupont and J. M. Tarascon, J. Electrochem. Soc., 2002, 149, A627.
17. Y. M. Lin, P. R. Abel, A. Heller and C. B. Mullins, J. Phys. Chem. Lett., 2011, 2, 2885.

---