New composite materials for lithium-ion batteries

Electrochimica Acta 84 (2012) 145–154

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New composite materials for lithium-ion batteries Brian L. Ellis, Kaitlin Town, Linda F. Nazar ∗ Department of Chemistry, University of Waterloo, 200 University Ave. West, Waterloo, Ontario, Canada N2L 3G1

a r t i c l e

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Article history: Received 25 February 2012 Received in revised form 28 April 2012 Accepted 30 April 2012 Available online 8 May 2012 Keywords: Li-ion battery Energy storage Polyanion positive electrode materials Silicon–carbon Negative electrode materials

a b s t r a c t Sparked primarily by the need for safe, portable, high-voltage energy storage, lithium ion (Li-ion) batteries have been heavily researched over the past three decades. The advancement of nanostructured electrode compounds has opened up the field of polyanionic materials and Si-based composites as positive and negative electrode materials respectively. Promising new materials with high capacity and high potential have driven interest in this field. This review will highlight and summarize developments of the field in the past 5 years, with a focus on polyanionic materials based on phosphates, silicates and sulfates and Si/C nanostructured compounds. Crown Copyright © 2012 Published by Elsevier Ltd. All rights reserved.

1. Introduction Global energy demand from both the grid and portable applications such as hybrid electric vehicles (HEVs) have created a need for environmentally responsible energy storage. Lithium-ion (Liion) batteries are one such energy storage system which have been investigated intensely owing to their high energy density, high operating voltage and low self-discharge. Current commercial Liion cells are composed of LiCoO2 and derivatives as the positive electrode and carbon-based materials at the negative electrode. First commercialized in 1991, LiCoO2 has several merits, most notably high electronic and ionic conductivity as a result of the material’s layered structure, although LiCoO2 /C cells are not suitable for all applications as a result of concerns over safety, cost and the environment. Efforts to improve and replace LiCoO2 /C cells have focused on the development of new positive electrode materials with minimal cobalt content such as polyanionic (phosphate, sulfate and silicate) compounds comprised of iron and manganese, which offer substantial safety and environmental benefits, as well as oxides of manganese and/or nickel. Of the many polyanionic phosphate, sulfate and silicate compounds, the most heavily researched are those in the olivine family, namely LiFePO4 and LiMnPO4 , first reported in 1997 by Goodenough and co-workers [1]. Other promising positive electrode

∗ Corresponding author. E-mail address: [email protected] (L.F. Nazar).

materials include fluorophosphates and fluorosulphates of the tavorite group including LiFePO4 F [2,3], LiVPO4 F [4,5], and LiFeSO4 F [6,7]; layered fluorophosphates such as Na2 FePO4 F [8]; and silicates such as Li2 FeSiO4 [9] and Li2 MnSiO4 [10]. The structures of these compounds are shown in Fig. 1. Each of these polyanionic materials suffers from poor electronic conductivity and much space in the literature has been devoted to conductive coatings to enhance particle surface conductivity and nanoscale synthesis to minimize the path length of lithium and electron migration in these materials, the highlights of which are summarized here. Other highly researched positive electrode materials include core/shell oxides such as Li[(Ni0.8 Co0.1 Mn0.1 )1−x (Ni0.5 Mn0.5 )x ]O2 , which have been summarized elsewhere [11]. Lithium–metal alloys, such as lithium–silicon, are among the most promising materials to replace current carbon-based anodes [12]. Silicon has a theoretical gravimetric capacity approximately 10 times greater than carbon (4200 mAh/g vs. 375 mAh/g). However, it is well-known that silicon and similar alloy materials undergo a large volume expansion up to 400% upon alloying with lithium over charge–discharge, which leads to a physical cracking and pulverization of the electrode and loss of electronic contact, poor capacity retention, and ultimately in cell failure. The focus of recent research within the field has been on solutions to overcome the aforementioned problems associated with volume change in Si-based anodes. Over the past 5 years, there have been many developments for both positive and negative battery materials, which will be summarized herein.

0013-4686/$ – see front matter. Crown Copyright © 2012 Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.electacta.2012.04.113

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Fig. 1. Graphical representation of key polyanionic structures: (a) the olivine structure of LiFePO4 and LiMnPO4 ; (b) the tavorite structure of LiVPO4 F, LiFePO4 F and LiFeSO4 F; (c) the layered fluorophosphates structure of Na2 FePO4 F and (d) the structure of Li2 FeSiO4 . The metal polyhedra are shown in blue, polyanion tetrahedral in yellow and alkali ions in grey. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of the article.)

2. Positive electrode materials 2.1. Olivines Lithium phospho-olivines were first reported as positive electrode materials for lithium-ion cells in 1997, when studies on LiFePO4 and Li(Fe,Mn)PO4 were published by the Goodenough group [1]. The olivine structure is shown in Fig. 1a. Transition metal octahedra (MO6 ) alternate edge and corner sharing with phosphate tetrahedra (PO4 ) to make a stable three-dimensional framework. Lithium ions reside in channels which run parallel to the b-axis within this structure. It is widely acknowledged that these b-axis tunnels are the primary route of Li transport in this structure: the results of AC impedance measurements on single crystals of LiFePO4 have shown that ionic conductivity is highest along the b-axis direction [13]. Structural simulations showed that the migration between adjacent Li sites in that direction did not occur by a linear hop, but rather along a curved pathway [14]. This result was confirmed by Nishimura et al. by performing maximum entropy method modeling on neutron diffraction data of a partially delithiated sample [15]. LiFePO4 has several advantages over other iron phosphate intercalation materials including the highest theoretical gravimetric capacity of any iron phosphate (170 mAh/g) and as a consequence of phosphate bonding in the structure, a high potential (3.5 V vs. Li) while the manganese olivine, LiMnPO4 , has a higher potential of 4.1 V vs. Li [1]. Malik and co-workers determined the functionality of LiFePO4 was the availability of a single-phase transformation path at very low overpotential [16]. Once a particle is partially lithiated or delitiated, charge or discharge continues with a small overpotential present and continues to insert or remove Li as rapidly as their diffusion and surface transfer kinetics allow. This allows the system to bypass nucleation and growth of a second phase, although the system will relax to produce a phase boundary once the overpotential is removed.

LiFePO4 has a low electronic conductivity of ∼10−9 S/cm [17] and like other insulating transition metal phosphates is considered to be a small polaron material. Long-range electrostatic interactions were calculated for a hole polaron and a lithium vacancy in LiFePO4 . The binding energy calculated for this interaction was 0.5 eV [18]. Experimental evidence of highly coupled lithium ion/electron transport was observed upon the formation of high temperature solid solution regimes of partially delithiated LiFePO4 : the onset of lithium disorder in these materials was determined to be 220 ◦ C [19], the same temperature at which rapid electron hopping was observed [20].

2.1.1. Synthesis and coating of olivine LiFePO4 and LiMnPO4 Synthetic methods for preparing olivines have focused on the production of both carbon coatings to improve surface electronic conductivity and nanoparticles to reduce the path length of lithium in the material in order to achieve good electrochemical performance of the material at high rates, as summarized in a previous review [21]. Alternate surface modifications have been found to enhance the ionic conductivity of the olivine particles: stoichiometries of the type LiM1−2x P1−x O4−y (M = Fe, Mn) [22,23] were prepared to produce an ionically conductive lithium phosphate glass on the particle surface which increased the electrochemical performance of both materials compared to uncoated LiFePO4 and LiMnPO4 . The capacity improvement is most noticeable in the iron compound: at a discharge rate of 2 C (2 Li per transition metal per hour), the discharge capacity of LiMn0.9 P0.95 O4−y is 95 mAh/g whereas for LiFe0.9 P0.95 O4−y , the full capacity (170 mAh/g) was obtained, as shown in Fig. 2. In a separate experiment, cells of LiFe0.9 P0.95 O4−y displayed appreciable discharge capacity at very fast rates (up to 400 C), although these cells were charged at a very slow rate (
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Fig. 2. Comparison of the discharge behavior of LiMn0.9 P0.95 O4−ı . Both materials were discharged at 2 C.

LiMn0.9 P0.95 O4−ı

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and

Reproduced with permission from [23]. Copyright 2010, The Electrochemical Society.

origin, identity and role of the surface coating on these materials became the subject of debate within the community [24]. Although the synthesis of nanoparticles alleviates the difficulty associated with ionic and electronic transport, nanoparticles tend to pack poorly, resulting in low tap density and poor volumetric performance of the cell. Sun and co-workers have found a new synthetic method for preparing nanoparticulate olivines which agglomerate into micron-scale particles in solution [25]. After annealing, the nanoparticles were carbon-coated and the agglomerates were found to be porous as a result of the decomposition of the glucose in the synthesis. SEM and TEM micrographs of a cross-section of one of these micron-sized agglomerates are shown in Fig. 3a and b. While electrodes of Li(Mn0.85 Fe0.15 )PO4 comprised of nanoparticles produced from spray pyrolysis displayed a gravimetric capacity of 165 mAh/g at C/20, electrodes prepared with the agglomerated nanoparticles exhibited a reduced capacity of 145 mAh/g. However, when considered on a volumetric basis, the agglomerated nanoparticles showed a clear advantage. The volumetric capacity of both the nanoparticles and micron-sized agglomerates of nanoparticles is shown in Fig. 3c. The agglomerates achieve a volumetric capacity of 370 mAh/cm3 and have a clear advantage in volumetric capacity owing to the high tap density of the agglomerates (1.4 g/cm3 ) compared to the standard nanoparticles (0.3 g/cm3 ). Clearly, both morphology and coatings must be considered when these polyanion materials are prepared. 2.1.2. Solid solutions, defects and nano effects in olivines The possibility of employing lattice modifications to introduce Li vacancies and inducing mixed-valence transition metals into the olivine structure of LiFePO4 and LiMnPO4 has intrigued researchers in this field for several years, even though initial reports of LiFePO4 described the electrochemical behavior as two-phase in nature [1]. Recent reports of doping manganese-rich olivines with small amounts (<25%) of iron, magnesium or both, have shown improved electrochemical performance, as compared to pure LiMnPO4 [25–28]. The iron dispersed in the structure act as nucleation sites for the delithiated phase, while the presence of magnesium reduces the lattice mismatch between lithiated and delithiated phases. Doping of vanadium on the M2 site of LiFePO4 was demonstrated by Omenya et al. who reported up to 10% V+3 could be substituted on the M2 sites for Fe+2 , thereby producing M2 vacancies [29]. Increased V doping improved the rate capability of

Fig. 3. (a) Cross-sectional SEM image and (b) cross-sectional TEM image of porous micro-Li(Mn0.85 Fe0.15 )PO4 . (c) Comparison of volumetric capacity of nanoLi(Mn0.85 Fe0.15 )PO4 and micro-Li(Mn0.85 Fe0.15 )PO4 . Reproduced with permission from [25]. Copyright 2011, John Wiley and Sons.

the material and close to theoretical capacities could be obtained at a rate of 1 C for the 10% V-doped material. The particle size of the as-prepared material defines several of the physical properties of LiFePO4 . The synthesis temperature greatly affects particle size and the formation of nanoparticles is often achieved at low temperatures, namely via precipitation from solution where the solvent acts as both a stabilizer and growth inhibitor for the particles. Nanoparticles of LiFePO4 containing significant lattice defects may be precipitated from water [30] whereas higher-boiling solvents such as ionic liquids [31] or long chain polyols [32–34], or sintered materials typically produce ordered olivines. A TEM study showed that bulk solid-state LiFePO4 prepared at 600 ◦ C contains approximately 1% antisite defects while LiFePO4 prepared at 800 ◦ C was virtually free of defects [35]. An early structural model published by Srinivasan and Newman predicted the existence of two single compositional regimes, Li1−y FePO4 and Lix FePO4 , close in composition to the end-members LiFePO4 and FePO4 [36]. These solid-solution regimes were first confirmed experimentally by Yamada et al. where X-ray diffraction studies on bulk particles of various compositions of LiFePO4 which had been chemically oxidized verified the presence of two phases: a Li-rich phase with a unit cell volume of 290.7 A˚ 3 , less than that typically found for pure LiFePO4 (291.1 A˚ 3 ) and a Li-deficient phase with a unit cell volume of 272.6 A˚ 3 , greater than that of pure FePO4 (271.9 A˚ 3 ). The limits of the solid solution were determined to be Li1−y FePO4 , y = 0.038 and Lix FePO4 , x = 0.032 using Vegard’s law [37]. A further study using neutron diffraction to determine lithium content showed the miscibility gap to be narrower, where y = 0.11 and x = 0.05 and these monophasic regions of the electrochemical curve strayed from the constant value of 3.42 V seen for the two-phase transition [38]. The first confirmed report of the direct synthesis of a mixedvalent olivine came in 2008 when Gibot and co-workers reported

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Fig. 4. Schematic representation of fit results for delithiation of nano-Li0.90 FePO4 which depicts the evolution of phase segregation of the Li-rich and Li-poor regions of the crystallites with regions free of Fe antisite defects delithiating before regions containing M1 site defects.

two such compositions, Li0.79 Fe0.97 PO4 and Li0.89 Fe0.96 PO4 , collected from aqueous precipitation [30]. The presence of extensive antisite defects and Fe3+ in each of these samples was confirmed. Both the electrochemical profile and in situ XRD measurements confirmed that cells prepared from the 40 nm material showed solid-solution behavior over the entire range of composition. More recently, Badi et al. varied the stoichiometry of the standard LiFePO4 polyol reaction to target Li-deficient stoichiometries [34]. Combined X-ray and neutron refinements confirmed that target stoichiometries of LiFePO4 , Li0.9 FePO4 and Li0.8 FePO4 produced [Li]M1 [Fe]M2 PO4 , [Li0.93 Fe0.01 ]M1 [Li0.01 Fe0.99 ]M2 PO4 , and [Li0.90 Fe0.03 ]M1 [Li0.03 Fe0.97 ]M2 PO4 respectively, whose particle size ranged from 40 to 100 nm in the longest crystallite dimension. A chemical delithiation study of [Li0.90 Fe0.03 ]M1 [Li0.03 Fe0.97 ]M2 PO4 was performed and the results are summarized in Fig. 4. Close examination of the diffraction data revealed the Li-deficient crystallites contained domains of localized defects. Delithiation to Li0.6 FePO4 revealed both a Li-rich and a Li-poor phase, each of which had an extended solid solution composition. While the Lirich phase retained the initial composition, the Li-poor phase had two key features: firstly, it was found to be free of antisite defects, indicating preferential delithiation of defect-free regions of the crystallites. Secondly, the composition of the Li-poor phase was determined to be Li0.14 FePO4 , a lithium content higher than any reported for any Li-poor triphylite phase formed by electrochemical methods. Further delithiation to Li0.5 FePO4 revealed that Li removed at this step was extracted from both the Li-poor phase and the Li-rich phase. Beyond this point, the defect-containing regions of the crystallite started to delithiate: delithiation to Li0.3 FePO4 revealed the presence of Fe on the M1 site in the Li-poor phase. Almost all of the lithium could be extracted chemically from this composition, even with the 3% Fe on the Li site, although the only 70% could be removed electrochemically at a slow rate. In both of the above studies, the Li-deficient olivines exhibited lower electrochemical capacity compared to pure LiFePO4 samples, owing to the Fe ions situated on the lithium sites which partially block Li tunnels in the structure. Ab initio calculations showed a decrease in the diffusion coefficient for Li transport by more than two orders of magnitude with defect concentrations as small as 0.5% [39]. A similar delithiation study was performed on hydrothermallyprepared LiMnPO4 by Chen and Richardson in 2009 [40]. Partial chemical delithiation of the manganese olivine revealed the lithium-rich phase (Li1−x MnPO4 ) does not exhibit a large stoichiometry variance as evidenced by the almost constant unit cell values observed at all the stoichiometries probed (see Fig. 5). In contrast, the Li-deficient phase (Liy MnPO4 ) formed showed significant range of composition: at 7% delithiation, the stoichiometry of the Li-deficient phase has a maximum lithium concentration of Li0.16 MnPO4 , as calculated from Vegard’s law from the unit cell

Fig. 5. Relationship between unit-cell volumes and domain size of (a) the Li-rich phase and (b) the delithiated phase in the x LiMnPO4 /(1−x) Liy MnPO4 two-phase mixtures, obtained from chemically delithiated LiMnPO4 . Reproduced with permission from [40]. Copyright 2009, The Electrochemical Society.

volumes shown in Fig. 5. As the degree of delithiation progresses, the lithium content of the Li-deficient phase decreases until full delithiation has occurred. Similar results were found with electrochemically delithiated samples of nano-LiMnPO4 as well. The miscibility gap between Lix FePO4 and Li1−y FePO4 has been found to reduce upon decreasing particle size: values of x = 0.06 and y = 0.12 are obtained for 40 nm particles [41]; values of x = 0.12 and y = 0.17 are obtained for 34 nm particles [42]; and values of x = 0.30 and y = 0.25 for 22 nm particles [43]. Further, Li1−y FePO4 solid solutions have been found after exposure of stoichiometric

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Fig. 7. Charge–discharge profiles recorded at C/20 rate at 55 ◦ C of Li2 FeSiO4 /C, where the Fe3+ /Fe4+ plateau at 4.8 V was accessed. Reproduced with permission from [64]. Copyright 2010, American Chemical Society.

Fig. 6. Charge–discharge curve of the layered Na2 FePO4 F cycled at C/10 vs. Li. Cycle 1 is shown in grey and cycle 2 in black.

to reintercalate back into the structure [51]. The material can also be cycled in a sodium battery and the average potential of these cells is 3.0 V [52]. Multifunctional materials such as Na2 FePO4 F are advantageous in terms of processing cost with regard to growing concern over lithium availability.

Reproduced with permission from [8]. Copyright 2007, Nature Publishing Group.

2.3. Silicates LiFePO4 nanoparticles, prepared at 400 ◦ C, to air [41]. Whereas bulk particles were found to be stable upon exposure to air, Yamada and co-workers found that 40 nm nanoparticles of LiFePO4 delithiated to Li0.93 FePO4 as indicated by the decrease in unit cell volume after air exposure. The lithium may form clusters on the surface of the nanoparticles. Simulations have shown that for particles which contain two phases separated by a phase boundary, the energy benefit of a phase boundary within a particle is reduced as particle size decreases to the point where the phase boundary is a significant fraction of the total particle size [44]. As a result, the miscibility gap will gradually narrow for particles of decreasing size. The electrode potential of nano-LiFePO4 also differs slightly from that reported in the bulk: nano-LiFePO4 has an equilibrium potential 5–10 mV higher than that observed for bulk LiFePO4 [42,45]. As a consequence, electrodes comprised of a mixture of bulk and nano-LiFePO4 were found to exhibit lithium transport from bulk particles to the nano particles upon equilibration after partial discharge [46]. 2.2. Fluorophosphates Fluorophosphates are another class of electrode materials which have been investigated as positive electrode materials for lithium-ion batteries. Depending on the connectivity of the ionic framework, these compounds may be expected to exhibit a high cell potential as a result of both the inductive effect of PO4 3− group and the electron-withdrawing character of the F− ion. One of the first successful fluorophosphate materials was LiVPO4 F, reported by Barker et al. [5,47,48], isostructural with the natural minerals tavorite (LiFePO4 OH) and amblygonite (LiAlPO4 F) whose structure is depicted in Fig. 1b. LiVPO4 F holds promise as a two-electron redox material: lithium extraction and insertion is based on the reversibility of the V+3 /V+4 redox couple at 4.2 V and the V+2 /V+3 redox couple at 1.75 V [5,49]. The electrochemistry of the iron variant of this group, LiFePO4 F, was first reported in 2010 and the Fe+2 /Fe+3 redox couple in this compound is at 2.75 V vs. Li/Li+ [2,50]. In 2007, a new sodium iron fluorophosphate (Na2 FePO4 F) with a layered structure was first reported [8]. The structure, shown in Fig. 1c, features corner-sharing of face-shared dimers of octahedral Fe+2 . When cycled vs. Li, the initial deintercalation of Na2 FePO4 F proceeds through a quasi-solid solution regime wherein Na2−x FePO4 F (x < 0.5) compositions are found to be single phase and the structure is maintained and the average potential of the cell is 3.5 V. The first cycle of this cell, cycled at a rate of C/10 is shown in Fig. 6. Upon discharge, lithium was predominantly found

As silicon is one of the most abundant elements in the Earth’s crust, silicates offer enormous potential for cost-effective positive electrode materials for Li-ion batteries and lithium metal silicates of the form Li2 MSiO4 (M = Mn, Fe, Co) have gained popularity as positive electrode materials. These compounds are all related to the various forms of Li3 PO4 : all the transition metal and lithium sites are tetrahedrally coordinated. The structures of the various silicate polymorphs have been summarized nicely in a recent review [53] and the structure of Li2 FeSiO4 in the P21 space group is shown in Fig. 1d. Owing to much similarity between these polymorphs, meticulous characterization has been carried out on some of these phases [54–60]. As with most silicates, these compounds exhibit very low electronic conductivity: 2 × 10−12 S/cm for Li2 FeSiO4 and 3 × 10−14 S/cm for Li2 MnSiO4 [61]. As such, one of the primary challenges with these materials is effective carbon coating to increase conductivity. Silicates have lower potentials compared to the corresponding transition metal olivines: the electrochemistry of Li2 FeSiO4 prepared at 750 ◦ C shows a plateau on the initial charge of 3.10 V vs. Li/Li+ , while on subsequent cycles this plateau shifts to 2.80 V [9,55,62]. This change in potential is indicative of a structural change during the first cycle [63]. The capacity of Li2 FeSiO4 is 166 mAh/g (extraction of only 1 Li), although a recent report from Manthiram and co-workers has shown that the Fe3+ /Fe4+ plateau, calculated to be at 4.8 V [10], may be reached electrochemically in cells cycled at 55 ◦ C [64], as shown in Fig. 7, Full lithium extraction would double the theoretical capacity of Li2 FeSiO4 to 332 mAh/g. Although it has a lower conductivity, Li2 MnSiO4 has been studied intently as a high capacity material with the expectation that the Mn3+ /Mn4+ plateau would be at a lower potential that that for the iron silicate [10]. Upon deintercalation of Li, Li2 MnSiO4 has been found to be amorphous [61] and in this state, quick fading of the electrochemical capacity was observed [59,61,65]. In spite of these drawbacks, the field of silicates is growing in popularity due chiefly to the large theoretical capacity of these materials. 2.4. Fluorosulfates In 2010, a new category of polyanionic materials suitable as electrode materials in Li-ion batteries emerged: lithium metal fluorosulfates. The initial compound studied in this class was LiFeSO4 F [6,7] which is isostructural with the mineral tavorite (see Fig. 1b). This material was initially synthesized by the solvothermal reaction of FeSO4 ·H2 O and LiF: the key to the reaction was thought

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to be the hydrophobic nature of the ionic liquid solvent, which moderated the water loss of the FeSO4 ·H2 O and permitted the topotactic conversion of FeSO4 ·H2 O to LiFeSO4 F [6]. Later, it was shown that LiFeSO4 F could be produced from the same reagents in hydrophilic solvents [7], polymers [66] and even a solid-state route [67], using identical precursors. LiFeSO4 F can also be produced by a non-topotactic method [68]. Regardless of the method of preparation, the initial reports of the electrochemistry of LiFeSO4 F were very positive: the voltage of the material (3.6 V) was found to be slightly higher than that of LiFePO4 and 80% of the 151 mAh/g theoretical capacity could be obtained at a rate of C/10, with good capacity retention displayed over 50 cycles; furthermore, 65% could be obtained at a C/2 rate [6,7]. One explanation to the successful electrochemistry is the relatively high ionic conductivity, determined to be about 4 × 10−6 S/cm near 150 ◦ C and roughly 3 orders of magnitude higher than that of the olivine at the same temperature [6]. The increased ionic conductivity is partially a result of the cavernous nature of the tavorite structure which has intersecting tunnels in three dimensions. Calculations of activation energies for Li transport through all the different tunnels gave values of 400–500 meV [69]. Mixed transition metal fluorosulfates with this structure could be synthesized (i.e. LiFe1−x Mx SO4 F, M = Mn, Co, Ni), however the higher potential of the substituted transition metals was too high to observe any electrochemical activity with standard organic electrolytes [70]. A key drawback with this compound is its relatively unstable nature: the compound is very hygroscopic will react with trace amounts of water in solvents used to wash the sample after synthesis to reform the precursors [67]. LiFeSO4 F has a relatively low thermal stability: decomposition commences at 350 ◦ C [7], and the material has the propensity to oxidize after ball-milling, which suggests nanoparticles of LiFeSO4 F are prone to oxidation [67]. A new series of fluorosulfates was recently reported with the same stoichiometry (LiFe1−x Mnx SO4 F, 0.1 ≤ x ≤ 1) but which adapt the triplite (Mn2 PO4 F) structure [71,72]. Owing to subtle changes in the structure, the voltage plateau for the iron redox couple for this compound is 3.9 V vs. Li/Li+ , the highest reported Fe2+ /Fe3+ plateau for any iron compound. A comparison of the electrochemical profiles of Li(Fe0.9 Mn0.1 )SO4 F with the tavorite structure and the triplite structure is presented in Fig. 8. In the case of triplite, the polarization was found to be quite large, even at low rates, indicative of poor transport kinetics, likely a result of Li/Fe/Mn mixing on the two cation sites [68,71]. As with the other fluorosulfate compounds, only the Fe plateau was accessible on electrochemical cycling. One of the key advantages of this material is the very small volume contraction upon delithiation: LiFeSO4 F with the tavorite structure experiences a lattice contraction of 10.1% on oxidation to FeSO4 F [7] while Li0.25 (Fe0.8 Mn0.2 )SO4 F prepared from the oxidation of triplite Li(Fe0.8 Mn0.2 )SO4 F experienced only a 0.6% reduction of unit cell volume [71].

3. Negative electrode materials To overcome the large volume changes in lithium–silicon alloys due to lithium insertion and extraction, the use of binder compounds, dispersion in active and inactive matrices and the effect of nano-scaled morphologies and thin films have been considered and summarized in a previous review of the field [73]. More recently, the use of Si/C compounds has dramatically increased capacity retention within these electrodes [74], and thus will be the main focus of this section. There are many different forms of Si/C composite anodes, with preparation methods including pyrolysis [75–77], sol–gel synthesis [78], mechanical milling [79,80] and chemical vapour deposition (CVD) [81–83]. As carbon has a lower surface area than Si and forms

Fig. 8. Charge and discharge curves for (a) tavorite and (b) triplite phases of the composition Li(Fe0.9 Mn0.1 )SO4 F discharged at a C/20 rate. The plateau of the redox potential is increased from 3.60 V in tavorite to 3.90 V in triplite for the same composition. Inset: the derived curves (−dx/dV) are plotted as a function of potential. Reproduced with permission from [71]. Copyright 2011, Nature Publishing Group.

a more stable solid electrolyte interphase (SEI), it produces more stable electrochemical cycling. Reversible charge capacity of about 1000 mAh/g has been observed with Si/C composites for hundreds of cycles, however gradual loss of contact between Si and C during discharge caused lowered coulombic efficiency and a decrease in charge capacity. This decreased capacity causes cell failure, and is the driving force behind current materials research [82]. 3.1. Si/C bulk Some early examinations of Si/C composites were conducted on materials synthesized by simple ball milling, which allows for easy control of homogeneity, composition, and particle size throughout the sample. Li insertion within a sample would cause the formation of Li–Si alloys, while the volume expansion of such compounds would be buffered by the more ductile graphite or graphite intercalation compound. Carbon-containing Si nanoparticles were prepared by ball milling of graphite and silicon mixtures [79]. In the ball-milled C/Si material (20% Si by mass), the first reversible capacity was 1039 mAh/g and a capacity of 794 mAh/g was maintained after 20 cycles. However, the initial irreversible capacity was greater than 100%. The initial coulombic efficiency could be improved by controlling the milling conditions. Nanocrystalline Si/mesocarbon microbeads (MCMB) composites prepared by ball milling for 5–20 h

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graphite particles. The initial discharge capacity of a 19.2 wt% Si anode was 1033 mAh/g composite, corresponding to a capacity of ∼4300 mAh/g Si (deviation from theoretical value due to SEI formation), and a coulombic efficiency of ∼80%. The reversible capacity decreased notably after 20 cycles, and failure occurred due to pulverization caused by volume expansion of Si particles exceeding the “cushion” of the C sol–gel matrix. Thus, while the impulse behind a bulk Si/C electrode material is valid, providing a conductive buffer with good dispersion, practical limitations of contact and extreme volume change indicate that bulk materials are unable to function effectively over multiple cycles. 3.2. Si/C nanowires

Fig. 9. Charge and discharge curves for (a) nano Si and (b) nano Si-MCMB electrodes. Reproduced with permission from [84]. Copyright 2004, Electrochemical and Solid State Letters.

were prepared as negative electrodes by Liu and co-workers [84]. The ball milling process was used to disperse nano-Si in the MCMB matrix that produced small local domains of Si surrounded by larger particles of MCMB. When the ball milling time was increased a better dispersion was obtained; however, the MCMB particles were broken. These broken particles caused irreversible reactions due to their increased surface area and the resultant formation of a passivating film. A consequential compromise between better dispersion and intact MCMB particles yielded the 10 h milled sample with the best performances such as a reversible capacity of 1066 mAh/g, an initial cycling efficiency of 90.7% and about 65% capacity retention after 25 cycles. While nanocrystalline Si generally exhibits a plateau during charge and discharge (Fig. 9a), examination of the charge–discharge curves of the composite material show the electrochemical behavior is largely due to alloying of Li–Si, indicated by a sloped curve (Fig. 9b). Thus the material is essentially a Si storage phase, surrounded by a minimal Li-storage C matrix. The stability of the material, with regards to cycling, is therefore highly dependant upon a uniform distribution of Si and high adhesion strength between the Si and C matrix. Thus an even dispersion would allow for good electronic contact throughout the material, and discourage aggregation of the Si zones. The motivation behind a bulk Si/C compound synthesized via sol–gel was similar to that of mechanical milling. However, the use of a sol–gel synthesis could introduce a further cushioned open three-dimensional structure, discussed by Niu and Lee [85]. In this work, a sol–gel graphite matrix was prepared from commercial TimCal graphite and methyltrimethoxysilane, then imbibed through mixing with Si (325 mesh). This resulted in an even distribution of Si over an open 3D ceramic network penetrated by

The concept of Si nanowires holds promise, as this morphology allows for relaxation of the strain that leads to pulverization in bulk electrodes, has a short diffusion distance to improve reaction kinetics, and can maintain direct contact with the current collector to discourage electronic isolation. The implementation of a carbon core in these nanowires would improve electron transport pathways and provide stable mechanical support, as C undergoes little structural or volumetric change upon cycling. Cui et al. [86] demonstrated the benefits of core–shell C/Si nanowires through the use of commercially available carbon nanowires, drop-cast on stainless steel current collectors, and coated via SiH4 CVD with a Si/C ratio of 3:1. When cycled against Li at a rate of C/5, a reversible capacity of ∼2000 mAh/g was obtained for the first 30 cycles, and a coulombic efficiency of 98–99.5% was obtained over 55 cycles. When cycled at a rate of 1 C, a capacity of 800 mAh/g was obtained, indicating good reaction kinetics due to the short path of the nanowires. However, while the nanowire structure was preserved over cycling, TEM images after 5 cycles indicate that the core–shell structure was lost, with no distinction between the C and Si zones, and nanopore formation in Si. Kim et al. [87] explored an alternative method for the formation of mesoporous Si/C nanowires, using a hard template of SBA-15. These nanowires contained 6% C by weight (CHS analysis), and Si nanocrystals of approximately 6.5 nm. A pore size of 2.3 nm and a pore-wall size of 6.5 nm were calculated from BJH analysis, and confirmed by TEM. Upon cycling at a rate of C/5, an initial charge capacity of 3163 mAh/g was obtained with a coulombic efficiency of 86%. An irreversibility of 14% was attributed to formation of the SEI, due to the large surface area of the nanowires. After 80 cycles, the capacity was reported as 2738 mAh/g, which corresponds to 87% capacity retention. While ordered mesopores initially provided a buffer for volume changes between nanowires, following 80 cycles some of the porous nanowires were pulverized while others aggregated to form nanoparticles (Fig. 10). This effect is similar to the observation of pure Si nanowires after 10 cycles [88]. This aggregation is likely due to non-uniform volume changes in the nanowires, and is supported by the exclusive presence of an amorphous phase within the aggregate. Thus reported core–shell nanowires appear unable to retain their structure, and therefore function, upon repeated cycling. 3.3. Si/C nanofiber Several studies on Si/carbon nanofiber (CNF) electrodes have been conducted by Si et al. [89,90]. The synthesized materials consisted of a carbon-coated nano-Si, with CNF added to provide improved contact between the electrode materials and the copper current collector. The Si/C composite was first produced from nano-Si powder and PVC (1:9 weight ratio), mixed in THF, dried, then heated under 2% H2 –Ar at 900 ◦ C for 2 h; the resulting product was then mixed with CNF and cast as an electrode. After 50 cycles, the composite electrode material showed no surface cracking,

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the low capacity retention prevents similar materials from commercialization. 3.4. Other Si/C morphologies

Fig. 10. TEM image of Si/C nanowire electrode after 80 cycles. Pulverization of the nanowires has occurred. Reproduced with permission from [86]. Copyright 2008, American Chemical Society.

however separation between the electrode and current collector was noted in one case of impedance studies. A low charge transfer resistance associated with lithium insertion and removal was reported for the composite further mixed with CNF, and attributed to good electronic contact between the active material and current collector due to the CNF [91]. The capacity retention was reported to be about 75% after 30 cycles. While an interesting proof of concept,

Conventional composites traditionally formed through pyrolysis or milling result in Si particles embedded in a C matrix, which cannot wholly buffer against the large volume change due to lithiation/delithiation. A porous composite would allow for volume changes, as well as facilitate fast transport of Li+ . The presence of carbon in the composite aids in the formation of the SEI and increases electronic conductivity. Magasinski et al. [92] produced a benchmark material following a bottom-up approach, using a liquid binder to trigger granulation, yielding particles of 15–30 ␮m (Fig. 11). Carbon black particles were annealed to give branched short chains, which were then covered with Si via CVD. A hydrocarbon, propylene, was used to initiate granulation, and prevent oxidation of the Si particles. The size of the carbon black branches and deposited Si particles determined the porosity of the granules. BET measurements of the specific surface area (SSA) was 24 m2 / g, and the pore size distribution ranged from 30 to 100 nm. The pore size was confirmed through SEM, and the highly porous structure provides space for stable volume expansion due to Li–Si alloying. The material was tested in coin cells against Li, reaching a reversible capacity of 1950 mAh/g at C/20. The specific capacity of the nanoparticles of Si alone was estimated to be 3670 mAh/g, the highest reported, with a coulombic efficiency of 100%. At discharge rates of 1 C and 8 C the capacity was 1590 and 870 mAh/g respectively. These relatively high rate capabilities indicate good reaction kinetics, further supported by a narrow, intense peak at 0.5 V in the differential capacity curve, associated with Li extraction. While not strictly within the prescribed range of this review, several developments within the past 4 years have brought pure Si anode materials into favor. Nanostructured Si materials that include a void to allow for volume expansion have recently prompted interest. Two materials selected to highlight this field are sealed Si nanotubes and interconnected Si spheres. While a promising geometry, with discharge capacities of approximately 3000 mAh/g over 10 cycles achieved, the use of

Fig. 11. Schematic of synthesis (a) and electrochemical performance (b) of Si/C granulated material. Reproduced with permission from [92]. Copyright 2010, Nature Publishing Group.

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repeated cycling indicates promise and provides a base for future work.

4. Conclusions The first report of LiFePO4 in 1997 sparked much research on polyanionic materials as positive electrode materials for Li-ion batteries. Despite a low electronic conductivity and the possibility of hetero-ions on the Li site disrupting ionic conductivity, intense research and debate on LiFePO4 within the community has been devoted to new synthetic methods, coatings, transition metal substitution and mechanism of (de)insertion. These findings have culminated in the recent commercialization of LiFePO4 -based cells. In the last 5 years, other groups of polyanionic materials have come to prominence: fluorophosphates and fluorosulfates hold great promise as high-potential electrodes and silicates (Li2 MSiO4 ) have drawn attention as high-capacity materials, although more work is required to make these new compounds commercially viable. From the first report of Si-based anode materials to present day studies, the main challenge to Si anodes is the large volume expansion due to lithiation, and all associated effects. From the variety of techniques that have been investigated to overcome this volume change, careful attention to the three-dimensional structure of the electrode is rapidly gaining popularity. The capacities achieved and retained from Si-based materials that include a void to accommodate volume changes due to lithiation/delithiation are currently the state of the art and will direct future research in the field. Studies to probe new compounds and frameworks suitable as electrode materials for Li-ion batteries with attributes such as high potential, capacity and long cycle life are an on-going challenge undertaken by the scientific community. Fig. 12. SEM images of interconnected hollow spheres (a) before and (b) after cycling. While the wall of the spheres thickens due to cycling, the particles remain intact. Reproduced with permission from [94]. Copyright 2011, American Chemical Society.

References

open-ended Si nanotubes shared the shortcoming of bulk Si, with a large capacity loss of almost 50% over 50 cycles [93]. The use of sealed nanotube arrangements provides space to accommodate the volume change associated with Li insertion/extraction, while limiting the surface area exposed to electrolyte and therefore SEI formation. Song et al. used a sacrificial template, followed by Si CVD and template removal via high temperature reduction. Polycrystalline Si nanotubes with an inner radius of 30 nm, thickness of 30 nm, and height of 3–5 ␮m were obtained. Another geometry evaluated of nanostructured Si was interconnected hollow spheres, by Yao et al. [94]. A uniform size distribution of this material was synthesized using a template approach, with SiH4 deposited via CVD on 350 nm silica particles; the particles were then removed with dilute HF. When cycled against Li at C/10, an initial discharge capacity of 2725 mAh/g was reached, and the capacity decreased 8% per 100 cycles. Ex situ TEM and SEM were employed to compare the pre- and post-cycled material (Fig. 12). An increase of the shell thickness from 25 to 46 nm, and the outer diameter from 200 to 237 nm indicated a 240% volume expansion. None of the spheres cracked, indicating reduced diffusion induced stress when compared with Si nanowires. While the surface of the spheres roughened after 40 cycles, the hollow center reduced exposure to electrolyte and therefore SEI formation. Additionally, the interconnected nature of this material reduces the need for binders and conductive additives. The capacities achieved and retained from Si-based materials that include a void to accommodate volume changes due to lithiation/delithiation are currently the state of the art and will direct future research in the field. A lack of materials degradation over

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