Ultimate Limits to Intercalation Reactions for Lithium Batteries

Review pubs.acs.org/CR

Ultimate Limits to Intercalation Reactions for Lithium Batteries M. Stanley Whittingham* NorthEast Center for Chemical Energy Storage, Binghamton University, 4400 Vestal Parkway East, Binghamton, New York 13902, United States systems. Alternatives include flywheels, superconducting magnetic energy storage (SMES), and batteries; only the last is technically and economically viable.1 This review will describe the present status of cathodes for Li batteries and build on the Chemical Reviews issue of 10 years ago while not duplicating that effort.2 The reader is referred there for a historical review of the field, as well as to the book Bottled Lightning.3 For materials not central to this review, the reader will be referred to several other recent reviews, for example, one on nonphosphate polyanions.4 However, a little review is CONTENTS appropriate to place what follows in context. The reader is referred to the 2004 paper by Winter and Brodd for an 1. Introduction 11414 2. Olivines 11415 excellent introduction to the electrochemistry basics of battery 2.1. Occurrence and Synthesis 11415 systems.5 2.1.1. Hydrothermal Synthesis 11416 The first rechargeable Li-ion batteries, using a TiS2 cathode6 2.1.2. Solid-State Synthesis 11417 and combined with a LiAl anode, were marketed by Exxon in 2.2. LiFePO4−FePO4 Phase Diagram and Defects 11419 the mid-1970s. These original cells are still operational today, as 2.3. Reaction Mechanism for LiFePO4−FePO4 indicated in Figure 1, which shows a paperweight containing a Intercalation 11419 TiS2 battery, a solar cell, and an liquid-crystal display (LCD) 2.3.1. Experimental Data 11420 watch. Ongoing tests show that these cells still maintain around 2.3.2. Models of Reaction Mechanism 11421 2/3 of their capacity, showing the intrinsic longevity of Li-ion 2.4. Substitution in the Olivine Lattice 11424 cells. It was not until 1991 that commercially successful Li-ion 2.4.1. Vanadium as a Case Study 11425 batteries were marketed by Sony using a C6Li anode and a 2.4.2. Other Substitutions 11427 7 LiCoO 2 cathode. Over the next decade Li batteries became 2.5. Stability of Olivine Materials 11428 dominant for powering portable electronics. However, the 2.6. Other Olivine Materials 11429 scarcity/cost of cobalt caused a search for alternatives where the 2.6.1. Lithium Manganese Phosphate 11429 cobalt is partially replaced by other metals such as in 2.6.2. Lithium Cobalt and Nickel Phosphates 11430 LiNi1/3Mn1/3Co1/3O2. As discussed later, essentially all these 3. Multielectron Phosphate Materials 11430 oxides can evolve oxygen under a high degree of oxidation, that 3.1. Vanadyl Phosphates 11431 is, high Li removal, and are thus intrinsically unstable under 3.2. Other Multielectron Phosphates 11435 4. Layered Oxides 11435 extreme use conditions. Then around 1996/1997 the electro4.1. Li- and Mn-Rich Layered Oxides 11436 chemical behavior of the phosphate LiFePO4 was reported.8 4.2. Thermal Stability of Layered Oxides 5. Conclusions and What Does the Future Hold Author Information Corresponding Author Notes Biography Acknowledgments Abbreviations and Specialized Terms References

11437 11439 11439 11439 11439 11439 11440 11440 11440

Figure 1. Timepiece powered by a Li-ion battery built around 1975, showing the LCD display, solar cell, and Li/TiS2 battery.

1. INTRODUCTION There is a growing demand for energy storage, for intermittent renewable energy such as solar and wind power, for transportation, and for the myriad portable electronic devices. By far the lowest-cost, large-scale energy storage is pumped hydro, but there are very few possible additional sites for such © 2014 American Chemical Society

Special Issue: 2014 Batteries Received: June 5, 2014 Published: October 29, 2014 11414

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Table 1. Comparison of Theoretical and Actual Storage Capabilities of Lithium-Ion Cellsa,b

a b

chemistry

size

Wh/L theoretical

Wh/L actual

%

Wh/kg theoretical

Wh/kg actual

%

LiFePO4 LiFePO4 LiMn2O4 LiCoO4 Si-LiMO4 Panasonic

54 208 16 650 26 700 18 650 18 650

1980 1980 2060 2950 2950

292 223 296 570 919

14.8 11.3 14.4 19.3 31.2

587 587 500 1000 1000

156 113 109 250 252

26.6 19.3 21.8 25.0 25.2

The theoretical values in the table assume only the active components, and no volume or weight for lithium beside that in the cathode. Reproduced with permission from ref 1. Copyright 2012 IEEE.

In this review, these phosphates will be first discussed and then compared with the well-studied transition metal oxides. These two systems have quite different electrochemistry; whereas the mixed metal oxides react mostly by a singlephase intercalation reaction, Lix[NiMnCo]O2, the electronically insulating olivine LiFePO4 was thought to react by a twophase reaction, LiFePO4 + FePO4. There have been many challenges with using a two-phase reaction and an electronically insulating solid as a battery electrode, which limited the initial capacities to only ∼0.6 Li.8 These challenges have now been overcome by several breakthroughs, including conductive coatings9 or a redox mediator,10 combined with nanosizing of the material. It was these breakthroughs plus the riveting paper by the Chiang group,11 suggesting, incorrectly as it turned out, that substitution had a major impact on the electronic conductivity, that got the scientific community activated. At the same time, the company A123 successfully commercialized LiFePO4 batteries for portable and stationary systems. However, despite these successes, the reaction mechanism is just now being fully understood, as is the role that defects and metal substitution play. The electrochemical behavior of natural triphylite LiFe1−yMnyPO4 has been determined9c and found to be under 80 mAh/g ( 0.21 Reprinted with permission from ref 21. Copyright 2009 American Chemical Society.

°C, there must be an alternate defect present to explain the larger unit cells; it is believed that this is the LiFe− + FeLi+ disorder defect, which on annealing can be removed giving the ordered phase. As noted previously, the lattice parameters of the unit cell depend on the cation ordering; thus, they can be used to determine whether an ordered, and as will be shown later electrochemically active, material has been formed. The structure of hydrothermally grown LiFePO4 single crystals was determined in 1977;22 the measured density was 3.564 g/ cm3 and the lattice parameters were a = 10.334 Å, b = 6.010 Å, c = 4.693 Å, and volume 291.47 Å3. Streltsov et al.23 redetermined the structure in 1993 and found almost identical values of 10.322 Å, b = 6.010, c = 4.692 Å, and volume = 291.4 Å3. This is the generally accepted value today irrespective of the synthesis method; where the volume found is outside the range 291.4 ± 0.2 Å3, the material almost certainly contains defects or the particles are very nano, < 50 nm. The iron-rich [Li0.0938Fe0.031]FePO412 of the previous paragraph has the following parameters: a = 10.345 Å, b = 6.003 Å, c = 4.696 Å, and volume 291.66 Å 3 , slightly larger than for the stoichiometric ordered material. In all of today’s studies, it is assumed that the excess iron is present as ferrous, not as ferric as proposed in some earlier work by Goni et al.24 for the phase Li0.7Fe0.1MgPO4; it should be noted that neither the reaction atmosphere nor the unit cell values were described. In a related study, they reported25 the formation of Li1−3xFexNiPO4 (0 < x < 0.15) in air at 800 °C, again proposing the presence of ferric ions on the Li site. In both cases the unit cell size increases with metal content on the lithium site, just as for the hydrothermal data shown in Figure 3; for these compounds and the iron-rich hydrothermal materials, the a and c lattice parameters increase while the b parameter decreases with increasing metal on the Li site. Studies of the reaction mechanism of the hydrothermal formation of LiFePO4 show at least one intermediate. Ellis et al.26 observed, when starting with an ammonium salt and LiOH, the formation of an amorphous green−yellow solid, which on heating first formed a phase similar to NH4FePO4· H2O, which then converted to LiFePO4. In contrast, when they added ascorbic acid to the reaction medium, crystalline yellow vivianite, Fe3(PO4)2·8H2O, was formed on mixing the reagents.26 Chen et al. designed an in situ autogenous pressure

methods. The latter include direct solid-state reactions and carbothermal processes, which allow for the ready use of the much lower cost ferric salts for the large-scale formation of LiFePO4. The manufacturing cost of this material is still a hurdle in its use for large battery systems. Thermal treatment is also a key step in producing an electronically conductive film on the LiFePO4. Synthesis temperature and environment are also critical to making electrochemically good material. 2.1.1. Hydrothermal Synthesis. The hydrothermal/ solvotermal synthesis of LiFePO4 has been the subject of much research and development over the past decade;17 it is now a commercial process in Candiac, Canada. The autogenous synthesis of LiFePO4 was first reported in 200118 but had rather poor electrochemical behavior. Yang et al. reported19 that such hydrothermally formed LiFePO4 has a significant concentration of Fe on the Li sites that reduces lithium diffusion in the tunnels and, hence, its reactivity. Subsequently, Chen and Whittingham reported20 that LiFePO4 formed hydrothermally had a lattice volume that decreased in size as the hydrothermal reaction temperature was increased, reaching a constant value at ∼291.3 Å3 at 200 °C and above, as shown in Figure 3. Rietveld analysis showed that this volume increase was due to iron on the lithium site, and this was associated with disorder of the type [Li−Fe + Fe+Li]. The plateau is an indication of ordering of the Li and Fe atoms on their own sites and has a value typical of materials formed at high temperatures. Zavalij and co-workers subsequently showed12 from single-crystal Xray diffraction that some hydrothermal LiFePO4 is in fact slightly Fe-rich, Li1‑nyFen+yFePO4, where y is ∼3%. Axmann et al.21 changed the Li/Fe ratio by varying the amount of Li3PO4 to Fe in the reaction medium. They found that the structure formed at 725 °C would not tolerate >1% Li vacancies on the Li site in Li1−yFePO4, and at lower lithium contents iron was found on the Li site, Li1−2yFeyFePO4. Even here when y exceeded 0.06, an impurity phase, sarcopside Fe1.5PO4, was identified by X-ray diffraction; 5% Fe1.5PO4 was reported for y = 0.06x, leading to an olivine composition of Li∼0.9Fe0.02FePO4. The unit cell volume is shown in Figure 2b and shows the same trends as reported for the hydrothermal data in Figure 2a, but with a much smaller change in volume. The differences in Figure 3 suggest that there may be several defects present, with the iron-rich compounds leading to unit cell expansions of 100 nm, two phases have been reported in a “single” crystal, e.g., Chen et al. using TEM56 and Laffont et al. using TEM/electron energy loss spectroscopy (EELS).57 This is clearly a metastable situation. (1a) In large crystals, extremely sluggish equilibration is observed whether performed in situ or ex situ, and using a solid polymer electrolyte such as poly(ethylene oxide) (PEO) compounds the issue, because there is no easy way for lithium ions to migrate from one crystallite to another.58 (1b) Solid solutions were suggested by Richardson and coworkers59 with a line phase of composition LiyFePO4, y = 0.60 11420

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Figure 11. Variation of lattice parameters as a function of Li concentration x in simferite, Li1−xMg0.5Fe0.3Mn0.4PO4. Reprinted with permission from ref 62. Copyright 2013 F. Omenya.

Figure 12. X-ray lattice parameters for highly defective LixFePO4. Reprinted with permission from 65. Copyright 2008 Macmillan Publishers Ltd.

reaction. Figure 14 shows the results of an in situ synchrotron X-ray diffraction study using a specially designed electrochemical cell.73 Figure 14a71 shows the impact of reaction rate for particles of 180 nm size, and the single-phase region increases as the reaction rate (C rate) increases. Figure 14b71 shows the impact of particle size and clear evidence for an increase of the single-phase region as the size decreases. These studies provide the most convincing direct evidence for the single-phase model even for particles as large as 180 nm. 2.3.2. Models of Reaction Mechanism. A number of models for the reaction mechanism have been proposed for the reaction of lithium with iron phosphate. For nanosize LiFePO4 it is becoming clear that the formation of an initial single phase, either Li1−αFePO4 or LiβFePO4, is critical to the reaction proceeding rapidly. This is equivalent to the need of a disordered lithium lattice. Once the lithium is disordered, then the reaction will proceed by a single phase so long as the ordering time is longer than the reaction time. Thus, a key to attaining a single-phase reaction is to attain an initial single phase; without that initial single phase, the new phase must be

Figure 13. Magnetic susceptibility of Li2/3FePO4. Reprinted with permission from ref 62. Copyright 2013 F. Omenya.

thermodynamic diagram of the two-phase system, resulting in continuous-phase transition during electrochemical reactions. (8a) Two recent studies, by Grey and co-workers71 and Wagemaker and co-workers,72 show the formation of a solid solution under high reaction rates consistent with a disordered lithium lattice that is unable to order in the time frame of the 11421

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Figure 14. In situ X-ray diffraction (XRD) shows clear direct evidence for a single-phase reaction for small particles where rate inhibits phase separation. (a) The solid solution region increases with the rate of reaction for 180 nm crystallites, and (b) the solid solution region increases as the particle size decreases. Reprinted with permission from ref 71. Copyright 2014 AAAS.

nucleated, e.g., FePO4 when lithium is removed from LiFePO4. The width of this initial single phase is a function of the crystallite size and the “perfectness” of the LiFePO4 crystal lattice; any lattice substitution will reduce the “perfectness” and

enable the disordering of the lithium ions. If the reaction is stopped partway, then the metastable phase LixFePO4 might split into the two thermodynamically stable phases by a spinodal decomposition mechanism.74 However, it is not clear 11422

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Figure 15. Domino-cascade model, showing (left) interface region and (right) reaction front moving through the crystal. Reprinted with permission from ref 76. Copyright 2008 Macmillan Publishers Ltd.

Figure 16. Cycling behavior of (a) Li/Al and (b) Li/FeF2,80 showing the nucleation overpotential on discharge. Reprinted with permission from ref 80. Copyright 2011 American Chemical Society.

will be seen in an equilibrated sample of many particles; essentially, once the new phase is created, the reaction goes to completion in that particle (i.e., phase growth is much faster than nucleation). The reaction front moves too fast to be observed. This is a nucleation−growth model. Although a sharp boundary between the two phases is proposed, there would seem to be no reason why the boundary could not be significant in size and be a function of the crystallite size. At one extreme the boundary could be almost as large as the crystallite for nanocrystals, and for micron-size crystals it could be very small relative to the crystallite size as proposed by Wagemaker et al.78 (2b) In the larger-particle nucleation−growth model, a nucleation energy (overpotential) is required to nucleate the new phase, and then the new phase propagates through the crystal. Two phases are observed in partially reacted crystals. This is the Chen/Laffont model discussed earlier and supported by Dedryvère et al.79 Such a nucleation energy has been observed in many electrochemical conversion reactions, such as Li + Al and Li + FeF280 (see Figure 16). However, no such obvious overpotentials have been directly observed in intercalation reactions, such as Li + FePO4. Weichert et al.81 reported that for mm size crystals there is an induction period

how such a mechanism might operate in a single crystal, without forming, for example, a dislocation between the two phases formed. The different models are described below. (1) The core−shell proposed by Padhi et al.8 is an isotropic model (whereas diffusion in LiFePO4 is very anisotropic), so it can only be considered at the macroscopic agglomerate level, as used by Srinivasan and Newman,75 not at the microscopic primary particle level. Laffont et al.57 showed in an ex situ HREELS study that, in particles of 100 × 200 nm and 50 nm thick, the center of the platelet is always FePO4, whereas the peripherary is LiFePO4 whether charging or discharging. They did not see any transition region between the phases. (2) In nucleation growth models the new phase is first nucleated and then grows through the crystal. The two models described are applicable for nanosize crystals and for micronsize crystals, respectively. (2a) The domino-cascade model was proposed by Delmas et al.76 and supported by Brunetti et al.77 (“We conclude that the domino-cascade model is confirmed experimentally, or at least its foundation, i.e., the fully lithiated or delithiated state of the particles after partial delithiation.”). This model, shown in Figure 15, assumes a transition region in each crystallite between the two phases. This model assumes only two phases 11423

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expected.53 Two compositions may even be observed in true single-phase systems in large single crystals.) There is as yet no direct observable evidence for this metastable single-phase model for LixFePO4, but it is the most plausible of all the proposed models for nanosize materials. There is now much supporting experimental data supporting this model, including the vanadium-substituted experiments of Omenya et al. and the thermal studies of Yoo and Kang,86 which are consistent with a low overpotential needed to generate a single-phase reaction route. Numerous papers argue for one model or the other based on observations made after equilibration of the electrode,87 so we cannot discern between the various models for the reaction path. NRL did nice piece of work that suggests either the Ceder or the nucleation−growth model is correct.88 None of these models satisfactorily answers the question: if you have one single crystal and its overall composition is Li0.5FePO4, can it exist and, if yes, what is its structure? Yes, is almost certainly the answer, but it is in a metastable state. Malik et al. reviewed the various mechanisms.89 It is still not clear, at least to this author, that, for two phases with a very small lattice mismatch, the interface between the two phases could not be a gradual change in lithium composition over several unit cells rather than an abrupt sharp interface. In conclusion, it is now quite clear that, for battery-grade LiFePO4, that is, material of less than ∼200 nm, the key to fast diffusion in these materials is disordered lithium within the FePO4 lattice. This disordering is initiated by a single-phase region beginning at x = 0 or x = 1 in LixFePO4. The width of these single-phase regions is determined by the particle size, by substitution in the lattice, and by temperature. The wider the single-phase regions, the narrower is the miscibility gap and the smaller is the lattice mismatch between the lithium-rich and lithium-poor materials. In addition, once the lithium disorder is generated, then if the time it takes to order is greater than the reaction time, phase separation will not occur and the system will stay disordered throughout the reaction. The olivines are by no means the only electrode system where there is a raging debate of 2-phase versus 1-phase. The titanium oxide spinel is one example, and the high voltage spinel, LiNi0.5Mn1.5O4, is another.90 Another case is that of VSe2,91 which at equilibrium has two redox plateaus for the formation of LiVSe2 but probably reacts by a single-phase mechanism; in contrast, however, the reaction of LiVSe2 to Li2VSe2 must go by a two-phase mechanism, because in the latter the Li occupy tetrahedral sites and in the former they occupy octahedral sites. The reaction is still fast, so a new phase reaction does not have to be slow, as we know from the reactions of, for example, Li with Al or Sn.

followed by growth of a reaction front through the crystal with much cracking. Lawrence Berkeley National Laboratory (LBNL) work suggests a nucleation−growth mechanism for micron-size crystals.56 Chueh et al. suggested82 that the reaction follows a nucleation limited process. (3) In the single-phase model the reaction proceeds from x = 1 to x = 0 in the single-phase LixFePO4, as in LixTiS2, but in contrast to LixTiS2 it is a metastable state, which on relaxation converts into two phases where each crystallite is one of the two phases. This model83 requires an overpotential (activation energy) to expand the single phase, as shown in Figure 17.

Figure 17. Plot showing the change in potential, first in the singlephase region, then the overpotential required to create a continuous single phase. Once single-phase is created, point A, then the reaction proceeds downhill until the final equilibrium single phase, point B, is attained. Adapted with permission from ref 83. Copyright 2011 Macmillan Publishers Ltd.

Once activated, the energy is all downhill just as in a nucleation−growth mechanism, so that the reaction goes to completion in that crystallite. This overpotential will be a function of the existing disorder in the crystal and, therefore, should be a function of the level of substitution.63 This model suggests that there is a potential gap between a crystal being charged and one being discharged of ∼30 mV, which is close to that observed by Dreyer et al.84 and may explain the “hysteresis” reported by Chiang and co-workers.85 This potential gap has been shown to be a function of the substitution level in LiFe1−3y/2VyPO4.63 (Just as in LixTiS2, during the reaction the composition across the crystallite may/ will not be a constant as is clearly shown in Figure 18; i.e., a transition zone is observed and a concentration gradient is

2.4. Substitution in the Olivine Lattice

Almost any divalent cation can be substituted for iron, with complete solid solution in most cases, including, for example, Mg, Mn, Ni, and Co; Vegard’s law is obeyed.12 In addition there is ample literature indicating that Fe-rich compounds are known;24 although it was thought that the iron is present as ferric, it is more likely that it is ferrous with vacancies on the lithium site to charge compensate. Theoretical calculations concluded that “on energetic grounds, LiFePO4 is not tolerant to aliovalent doping (e.g., Al, Ga, Zr, Ti, Nb, Ta) on either Li (M1) or Fe (M2) sites.”54 Although it was not thought to be feasible to substitute more than ∼1% of an aliovalent ion on the Li site,92 it is now clear

Figure 18. Optical view of the lithiation of a TiS2 single crystal, showing the transition zone between LiTiS2 and TiS2. Reprinted with permission from ref 53. Copyright 1979 Elsevier. 11424

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oxygen, so it might substitute on the octahedral Fe site, forming a VO6 group, or on the tetrahedral P site, forming a VO43− group. Although there have been several reports on vanadium substitution to improve the electrochemical performance of LiFePO4, there is much disagreement about its role, varying from structure modification by substitution for Fe, P,106 or Li leading to enhanced Li or electron mobility, second phase formation, or formation of a conductive coating on the LiFePO4. In addition, there is no consistency in the reported lattice parameters and vanadium location in the lattice. Hua et al.107 and Ma et al.108 reported a linear decrease of the cell volume upon substitution of 4 and 7 mol % vanadium at the Fe site and the formation of Li3V2(PO4)3, while Sun et al.109 observed a cell volume increase with 3 mol % substitution and Zhang et al.,110 substituting up to 5% vanadium, observed irregular cell volume changes. Jin et al.111 observed no cell volume change upon adding up to 10 mol % V to the 1:1:1 Li/ Fe/P precursor ratio, concluded no vanadium substitution, and attributed the enhanced electrochemical performance to the conductive V2O3 incorporated into the carbon coating. Aleees112 reported the formation of a composite LiFePO4· zV2O3, with no V substitution in the LiFePO4 lattice but surprisingly showing a cell volume decrease to 288.9 Å. Zhao et al.113 claimed that “both experimental and theoretical simulations show that vanadium does not enter into the LiFePO4 crystal lattice”, not even at the 1% substitution level, based on the X-ray absorption data analysis. Harrison and Manthiram114 proposed that vanadium may be substituted into LiFePO4 as the vanadyl (VO)2+ ion. There is confusion about not only where the substitution resides but also what is its oxidation state, with reports of V3+,109 between 3 and 4+,110 or 4+;108 5+ is almost certainly not an option in the presence of Fe2+. Omenya and co-workers62,63,94 have made a systematic study of the substitution of part of the iron and/or lithium by vanadium. Vanadium was chosen not only to understand the discrepancies described above but also because of the large difference in its X-ray and neutron scattering factor, which allows for the ready determination of where all the ions reside in the structure. One key in the study was in the synthesis to tailor the precursor ratios for various oxidation states and charge-compensation mechanisms. This study clearly showed that vanadium can be substituted for some of either iron or lithium in LiFePO4. Figure 19 shows the change of lattice parameters for each case, clearly differentiating between the two cases. For substitution on the vanadium site, the volume decreases with the increase of vanadium,93 whereas for Li substitution the volume increases.94 In the latter case, the vanadium actually resides on the Fe site, displacing an equal amount of Fe to the Li site; charge compensation in this case is by Li vacancies, whereas for Fe substitution charge compensation is by Fe vacancies. The data in this figure is for samples synthesized at 550 °C. For the case of 20% V on Fe, a NASICON-type phase is formed; this NASICON phase is also formed at lower vanadium contents when the synthesis is performed at 700 °C. The three characteristic redox peaks for this phase are seen in the electrochemistry of these materials formed at higher temperatures. Such peaks are completely absent in the 10% V on Fe site materials formed at 550 °C. Harrison et al.98 also concluded that vanadium can be incorporated onto the iron site and that the solubility is very much temperature-dependent; for samples prepared by micro-

that substitution is quite possible on both the Li and Fe sites but that it is strongly temperature-dependent, with the solubility decreasing with increasing temperature.93,94 Such substitution also enhances the rate capability of the material.63 Meethong et al.95 also showed that substitution is quite possible for aliovalent cations; they suggested that ions like Zr4+, Nb5+, and Al3+ can be substituted on either the Li or Fe site but showed no definitive evidence for the substituted site preference; they suggested that charge compensation was achieved by Li vacancies. They did not consider the possibility that the aliovalent cation is on the Fe site with vacancies on the Li site. Delacourt et al.96 found no evidence in their studies for niobium substitution in LiFePO4 . Wagemaker et al. 97 questioned their level of aliovalent substitution in the olivine lattice, capping it at ∼3% for samples synthesized at 600 °C [it is now known that the solubility is a strong function of temperature].93 At that time there were numerous experimental and theory99 papers suggesting that it was impossible to have aliovalent substitution in LiFePO4; however, the paper of Herle et al.100 was not one of them, despite references51 to the contrary. It should be stated at the onset that there is no convincing evidence today that substitution has any significant effect on the electronic conductivity of the olivine, despite early reports to the contrary11,101 and subsequent discussions.102 Herle et al. showed100 that high-temperature annealing can cause the formation of conductive carbon and/or iron phosphide surface coatings; this has been confirmed by others. 103 The conductivity of the composite can also be improved by intentional formation of conductive carbon coatings or the addition of materials such as graphene.104 Appropriate substitution can be expected to have a major effect on the ionic conductivity of the lithium ions in the FePO4 lattice. With the appropriate defect the material could be converted from a one-dimensional conductor to a twodimensional conductor. For a Li ion to jump from one tunnel to an adjacent one, it must jump through an iron site, creating a defect pair LiFe + FeLi, the same one already discussed. This will cost a lot of energy so it is very unlikely in defect-free materials, but if the defect is already present the relevant activation energy is just that for the hop and should be quite feasible. This is exactly as in the analogous sodium beta aluminas, Na1+xAl1O17+d, where sodium ions move very rapidly with an activation energy of only 4 kcal/mol; in contrast, in stoichiometric NaAl11O17, there is essentially no sodium motion. The lithium ionic conductivity has been determined in large single crystals,105 where there is not expected to be any significant single-phase regions; the results indicated 2-D ionic conductivity. Much of the confusion over aliovalent substitution can be associated with two factors. First the solubility of aliovalent substituents appears to be a strong function of synthesis temperature. Second, it is essential to mix the proportions of the reactants to take the charge compensation in mind; if not, impurities are found, leading to much confusion in the literature. By the appropriate choice of synthesis temperature and reactant ratios, substitution seems to be possible on both the M1 and M2 sites. The next section will describe the case of vanadium substitution into LiFePO4. 2.4.1. Vanadium as a Case Study. Vanadium is a particularly attractive element to consider for substitution in the olivine lattice, as it readily forms phosphates and exists over a wide range of oxidation states and in different polyhedra with 11425

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Figure 20. Electrochemical behavior of LiFe1−3y/2VyPO4. (a) Charge− discharge curves measured at C/10 for the 550 °C synthesized materials and (b) cycling performance measured from 0.1 to 10C rate within the voltage window of 2.0−4.3 V.93 Reprinted with permission from ref 93. Copyright 2011 American Chemical Society.

Figure 19. Lattice parameters for vanadium substitution (black) LiFe1−1.5yVyPO4 and (red) Li1−3yFey[Fe1−yVy]PO4.93 Reprinted with permission from ref 93. Copyright 2011 American Chemical Society.

lattice. An in situ X-ray synchrotron diffraction study63 showed direct evidence that the phase diagram is significantly modified on vanadium substitution. Figure 21 shows that the single-

wave−hydrothermal at 280 °C, as much as 25% V was found in the structure. The electrochemical behavior of the LiFe 1−1.5y V y PO 4 materials formed at 550 °C is shown in Figure 20a.93 The capacity and cycling curves are almost identical for the 5% and 10% vanadium materials at this low C/10 rate. The 20% material shows several differences. First it shows three redox plateaus above 3.5 V; these are characteristic of the NASICON phase, Li3V2(PO4)3. This phase might be a mixed FeV phase. Second, there is a marked loss of capacity, which could be associated with the reduced amount of the olivine phase. The kinetics of the lithium intercalation and removal are also improved by the addition of vanadium on the iron site as indicated in Figure 20b. The capacity is the same at the end of cycling as at the beginning, as shown by the closing C/10 capacities. The capacity as expected decreases with increasing rate of reaction. Addition of 10% vanadium markedly improves capacity retention at all rates. However, most important of all is increasing the synthesis temperature to 700 °C, where much enhanced capacity retention is observed. This can be associated with several factors. First, conductive films are likely formed on the surface from both the formation of electronically conductive Fe2P impurities at these temperatures and the conversion of any residual carbon (from the synthesis precursors) from the insulating sp3 form to the conductive sp2 form, which occurs above ∼650 °C. The second factor is the presence of the electrochemically active Li3[VFe]2(PO4)3 phase. Hong et al.115 also showed the benefit of including activated carbons in the synthesis process for 5% V substitution in LiFePO4. The question remains as to why vanadium, or any other substitution, aids the intercalation of lithium into the olivine

Figure 21. Phase diagram for LiFe0.85V0.1PO4 compared with that of Dodd et al.46 for LiFePO4 showing the expanded single-phase regions. Data derived from refs 46 and 62.

phase regions are increased to ∼15% at room temperature, so that the miscibility gap is reduced to 1 and x is [Fe]. The stability window depends on a number of factors including temperature, particle size, presence of moisture, oxidizing or reducing environment, and carbon on the surface, among others. It is important to know when these compounds might decompose, particularly if this occurs with release of heat, which could exacerbate a thermal runaway situation in an operational battery. A systematic synchrotron X-ray diffraction study130 of the thermal stability of delithiated 100−200 nm olivine Fe1−yMnyPO4 in helium and oxygen atmospheres shows that the decomposition route is different. o-FePO4 loses oxygen to form Fe7(PO4)6 under inert conditions above 600 °C and to form single-phase t-FePO4 at temperatures above 600 °C in O2. Low manganese content materials, o-Fe0.8Mn0.2PO4 and oFe0.6Mn0.4PO4, both show structural stability up to 600 °C under inert atmospheric conditions. Interestingly samples containing Mn do not show any tendency to form trigonal, tFe1−yMnyPO4, phases. Between 600 and 700 °C, small amounts of sarcopside are found, and above 700 °C the pyrophosphate phase, (Fe,Mn)2P2O7, is formed. At higher manganese contents, Mn ≥ 0.6, some structural changes are observed at temperatures as low as 250−300 °C; these are still olivinerelated. o-Fe0.4Mn0.6PO4 shows sarcopside peaks between 550 and 700 °C, with pyrophosphate first appearing at 650 °C. For Fe0.2Mn0.8PO4 and Fe0.1Mn0.9PO4, the sarcopside phase appears

2.6. Other Olivine Materials

2.6.1. Lithium Manganese Phosphate. There has been increased interest in the behavior of LiMnPO4 because of its higher 4 V redox potential. It has significantly greater challenges than LiFePO4 as a battery cathode because its electronic conductivity is at least 4 orders of magnitude lower;132 there is a higher lattice mismatch between LiMnPO4 and MnPO4 than for the iron analogue, and there are lattice distortions when the Mn3+ ion is formed on lithium removal. Yamada and Chung133 showed that there is a complete solid solution between Fe and Mn in LiMn1−yFeyPO4 and noted difficulties in cycling these materials when y ≤ 0.2. Clément et al.134 confirmed that the Mn and Fe ions mix freely in LiMn1−yFeyPO4 for y = 0.25, 0.5, and 0.75 using advanced NMR techniques that overcome the complexities of paramagnetic materials, combined with density functional theory (DFT) studies. Piper et al.135 reported that the distortion of the MnO6 octahedron is a preferential elongation of two of the equatorial Mn−O bonds; the associated strong electron/lattice interaction is a key obstacle to the reversible lithium intercalation in pure LiMnPO4.133 Thus, it was very difficult to cycle most of the lithium until HPL showed the feasibility,136 and Drezen et al.137 showed the criticality of using very finely divided particles. They used a sol−gel synthesis method, and after calcination between 520 and 600 °C formed primary particles of size 140−220 nm. Subsequent ball-milling reduced the size to 90−130 nm 11429

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complex LiCoPO4, for its even higher redox potential of 4.9 V. The first electrochemical studies of cathodes containing LiCoPO4 were done in a study to increase the potential of LiCoO2 cathodes by the addition of phosphorus, which showed that the potential could be raised to 4.9 V.13 These were probably the first reported studies of an olivine cathode. Several early studies showed capacities ranging from only 70 mAh/g when charged to 5.3 V143 to exceeding 100 mAh/g,144 and Nakayama et al.145 have shown the reversibility up to 120 mAh/g when charged to 5.1 V; the latter also showed two plateaus on charging. The capacity can attain 120 mAh/g at 0.1 mA/cm2 by the addition of excess lithium in the form of Li3PO4 and Co3O4,146 but the capacity fell to half the value over 30 cycles. Surface coatings also allow the attainment of 120 mAh/ g.147 A carbon LiCoPO4 composite attained a capacity of >140 mAh/g at C/10, but just as with the examples above, the capacity dropped by 50% over 30 cycles;148 it showed high rate behavior with a capacity of 70 mAh/g at 20C. Just as with LiMnPO4 the substitution of part of the Co by Fe enhanced the capacity,149 but there is a large excess charge on each cycle and so the true characteristics of LiCoPO4 will only be determined when a stable electrolyte is found. It is now clear that, on removal of lithium from LiCoPO4, the phase Li0.7CoPO4 is formed first and then CoPO4. The lattice parameters of these three olivine phases have been determined by Bramnik et al.67 and are shown in Table 3 below. Whereas this intermediate composition appears to be thermodynamically stable for Co, it is only seldom reported for the iron system. The 70 mAh/g results reported are almost certainly due to charging only to the intermediate phase, based on the lattice parameters of the charged material (Table 3). The CoPO4 phase turns amorphous on exposure to air67,150 and appears to be even more sensitive to moisture than MnPO4. Antisite defects are also reported for LiCoPO4,151 just as found for LiFePO4. Although LiCoPO4 is still of interest, the stability of the electrolyte at the >5 V charging potentials has limited research. However, it has been evaluated against a titanium oxide spinel anode and gave capacities of ∼120 mA/g.152 The material has also been used as a coating on other cathode materials. Lee et al.153 reported that a LiCoPO4 coating on LiCoO2 significantly improves the latter’s stability to swelling and allows for capacities over 190 mAh/g at C/10. Kang and Thackeray154 also used a LiNiPO4 treatment to stabilize the capacity and increase the rate capability of the Li/Mn-rich layered oxide cathode materials.

depending on time of ball-milling; these particles had reversible capacities of 156 mAh/g at C/100 and 134 mAh/g at C/10. These are 92% and 79% of the theoretical capacities, respectively. Earlier results were mostly obtained at lower rates, for example, Delacourt et al.138 found capacities of 70 mAh/g for 100 nm diameter particles compared with only 35 mAh/g for 1 μm particles, both at C/20 rates. Yonemura et al.139 attained 150 mAh/g for small particles at C/100 rate. The HPL work reported that partial isovalent substitution of some of the manganese particularly by iron significantly improved the rate capability of the material as shown in Figure 26a.140 When

Figure 26. Capacity of (a) LiMn0.9M0.1PO4 as a function of substituent and rate140 and (b) LiMnyFe1−yPO4 as a function of Mn content.141 (a) Reprinted with permission from ref 140. Copyright 2010 The Electrochemical Society. (b) Reprinted with permission from ref 141. Copyright 2014 American Chemical Society.

3. MULTIELECTRON PHOSPHATE MATERIALS The maximum theoretical energy storable in a LiFePO4 cell is 587 Wh/kg and 2 kWh/L, and in a LiMnPO4 cell it is 684 Wh/ kg and 2.4 kWh/L, showing that these cells store too little energy to be viable for many portable applications, including electronic devices and probably personal automobiles. To increase the storage, a second electron must be incorporated per metal redox center. This might be accomplished by inserting two lithium ions or one magnesium ion. A number of cathode materials have shown the capability of storing two lithium ions by an intercalation mechanism. These include the layered dichalcogenides such as VSe2,91 the layered oxides such as LiNi0.5Mn0.5O2, and vanadyl phosphate, VOPO4.155 Only the last has both redox potentials at useful potentials for storing energy. Recently a database has been constructed of the expected potentials in a lithium cell for metal phosphates,156

the iron composition is increased, Ravnsbæk et al.141 showed in 2014 that the rate capability is further improved and, at the highest rates, is even better than that for pure LiFePO4, consistent with the substitution studies discussed previously. These higher iron content samples, such as LiMn0.4Fe0.6PO4, have been found141 to enable the metastable single-phase reaction mechanism, as discussed for the iron analogue. In this case, for particles below 150 nm, a continuous change in lattice parameter was found using operando, high-precision synchrotron radiation powder X-ray diffraction (SR-PXD) as shown in Figure 27 for two particle sizes, 52 and 106 nm. 2.6.2. Lithium Cobalt and Nickel Phosphates. More recently there has even been interest in the much more 11430

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Figure 27. Unit cell volume changes during Li extraction (charge) and insertion (discharge) at a C/10 rate for particle sizes of 52 and 106 nm LiMn0.4Fe0.6PO4. The vertical lines are the boundaries between the one- and two-phase regions reported by Yamada and co-workers48c (black lines) and Malik et al.142 (red line). (i), (iii), and (iv) denote the two-phase regions, while (ii) and (v) denote the one-phase regions. The smaller figures at the top and bottom show the evolution of the (200) peak within the two-phase regions during charge (upper) and discharge (lower).141 Reprinted with permission from ref 141. Copyright 2014 American Chemical Society.

Table 3. Lattice Parameters of LixCoPO4 Phasesa

a

3

compound

a, Å

b, Å

c, Å

vol., Å

ref

LiCoPO4 Li0.7CoPO4 CoPO4 LixCoPO4 LiCoPO4 LiCoPO4

10.1955 10.070 9.567 10.089 10.202 10.215

5.9198 5.851 5.7806 5.853 5.922 5.918

4.6971 4.717 4.7636 4.719 4.699 4.706

283.49 277.94 263.43 278.66 283.90 284.49

67 67 67 141 141 12

readily accessed. In addition, couples based on iron and manganese 4+/2+ are also of interest but fall above the red line, exceeding the electrolyte stability limit. The copper couple Cu1+/2+ looks particularly interesting, but it is likely that a displacement reaction, in which copper metal is precipitated out, will be preferred over lithium intercalation. 3.1. Vanadyl Phosphates

Vanadyl phosphate, VOPO4, is found in at least seven different crystallographic forms from layered structures to 3-dimensional lattices. All contain a short VO bond, varying from 1.6 to 1.8 Å, and a long V−O bond, around 2.2 Å. Confusion abounds about the nomenclature used to describe these different phases because the Greek symbols used were assigned in order of reporting of the phases. Thus, in the literature α-LiVOPO4 is not the lithiated form of the layered structure α-VOPO4 but rather that of ε-VOPO4. For clarity in this review, we will call this lithiated form ε-VOPO4 and similarly for the other phases, i.e., the delithiated name will determine the name of the

Republished with permission from ref 12. Copyright 2008 Elsevier.

among others. Figure 28 shows the calculated potentials for the metal phosphates. It can be immediately seen that there are not many options open if the cell potential is to be maintained below that of the decomposition potential of the electrolyte, < 5V. Two couples are circled, vanadium and molybdenum. The former is well-known to react with at least two lithium ions, and in oxides, such as V2O5, the valence states from 5+ to 3+ can be 11431

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Figure 28. Computational prediction of the cell voltages for a wide range of transition metal phosphates. Two of the most promising couples are identified in the red ovals within the electrolyte stability limit of 4.5 V (red line). Adapted with permission from ref 156. Copyright 2011 American Chemical Society.

Figure 29. Relationship between the various structures in the VOPO4 system; the building block for all these structures is shown at the lower right.

formed170 with a very closely related but triclinically distorted lattice with a slightly increased volume of 85.51 Å3. This structure was recently confirmed,160 and the chain-linking is shown in Figure 30. It is very similar to that of LiVPO4F, but in the fluoride the bonds along the chain have the same length. The structure of Li2VOPO4 has now been worked out.171

lithiated forms. Their structural parameters are shown in Table 4, and the fundamental building block and some relationships between the structures in Figure 29. The β-phase is the most dense of the lithiated phases so it might be expected to be the stable one at high temperatures. It has been prepared by carbothermal methods at 500 °C.157 The α-VOPO4 is the most dense of the VOPO4 phases, whereas its lithiated form αLiVOPO4 has almost the lowest density. Whether these structures can change phase on lithium intercalation and removal is not known but may explain some of the complex electrochemistry observed. Studies are underway to better understand the phase stability and lithium diffusivity.158 Lim et al.161 first reported in 1996 the existence of the εVOPO4 phase, which was formed by the thermal decomposition at 550 °C of H2VOPO4; heating at 700 °C formed the β-phase. The dehydrogenation reaction could be reversed by hydrogen spillover at 150 °C using a Pt catalyst. Song et al.155 showed that this hydrogen deintercalation reaction could also be performed in an electrochemical cell at room temperature, showing the high mobility of the protons. The structure of εVOPO4 is made up of chains of VO6 octahedra joined at opposite apexes;169 this O−VO link contains a short VO bond and a long V−O bond. There are two different vanadium sites alternating along the chain. A recent synchrotron study confirmed this structure, with a formula volume of 82.14 Å3.159a On lithiation, the previously known structure of ε-LiVOPO4 is

Figure 30. Comparison of [VO4X2] chains in LiVPO4X (X = F, O). Reprinted with permission from ref 160. Copyright 2012 American Chemical Society.

Table 4. X-ray Lattice Parameters of VOPO4 and LiVOPO4 phase

a, Å

b, Å

c, Å

β, deg

V, Å3

V/VOPO4

ε-VOPO4 ε-LiVOPO4160 ε-Li2VOPO4167 β-VOPO4161 β-LiVOPO4162 β-LiVOPO4163 δ-VOPO4164 ω-VOPO4165 γ-VOPO4165 α-VOPO4166 α-VOPO4167 α-LiVOPO4168

7.279 6.732 7.199 7.786 7.446 14.887 9.055 4.855 17.397 6.200 6.014 6.291

6.886 7.194 7.101 6.130 6.278 12.575 9.055 4.855 8.820 6.200 6.014 6.291

7.265 7.920 7.777 6.968 7.165 7.174 8.608 8.430 4.908 4.110 4.434 4.445

115.39 tricl tricl 90.0 90.0 90.0 90.0 90 90 90 90 90

328.57 342.03 356.36 332.57 334.94 1343.00 705.74 198.72 753.06 157.99 160.37 175.93

82.14 85.51 89.09 83.14 83.74 83.94 88.22 99.36 94.13 79.00 80.35 87.97

159a

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All the phases listed in Table 4 have been evaluated in electrochemical cells with varying degrees of success.172 Initial electrochemical data on ε-VOPO4 was reported by Kerr et al.173 and showed that one lithium insertion at ∼4 V to form LiVOPO4 was very reversible. Song et al. later started by first deintercalating the protons from H2VOPO4 in a lithium anode electrochemical cell to form the ε-VOPO4 phase at ambient temperatures,155 and then inserting one lithium at ∼4 V and a second at ∼2.5 V. Recently there has been a resurgence of interest in this material, with cycling studies and fundamental structural efforts underway. Figure 31 shows the cycling behavior and rate

bypassing the LiVOPO4 phase; this results in a significant loss in storage efficiency. Also apparent from Figure 31 are several redox features between 2.6 and 2.0 V for the material from the monoclinic H2VOPO4. These are even more apparent from the cyclic voltammograms shown in Figure 32a. These suggest

Figure 32. CV curves for (a) ε-VOPO4 from the tetragonal and monoclinic H2VOPO4 at a scan rate of 0.05 mV/s and (b) that from the monoclinic phase over a range of scan rates. Reprinted with permission from ref 159a. Copyright 2013 The Electrochemical Society.

there are three different lithium sites for the second lithium; they are much more pronounced for the material from the monoclinic precursor. Running the cyclic voltammograms over a range of sweep rates, as shown in Figure 32b, allowed a diffusion coefficient of ∼10−10 cm2/s to be calculated.159b Bianchini et al.171 have recently synthesized the three phases Li1.5VOPO4, Li1.75VOPO4, and Li2VOPO4 and determined their structures. That of the lithium-richest phase is shown in Figure 33, and the five different Li sites are identified. (In the figure the authors have used the tavorite-type notation of VPO4O, rather than that used in this paper of VOPO4 which emphasizes the VO bond.) As the lithium content increases, the V−O distances become more equivalent but not all the same as found in LiVPO4F, which is shown in Figure 30. Bianchini et al.171 synthesized their vanadyl phosphate as the lithium compound LiVOPO4 by a ceramic route resulting in crystallite sizes of ∼1 μm. As a result, in order to obtain welldefined redox plateaus, they milled the material for 30 min to reduce the size. The resulting very slow, almost open-circuit scans well represent the thermodynamic conditions and are shown in Figure 34. The differential plots clearly show the presence of the three compositions at x = 1.5, 1.75, and 2.0 in LixVOPO4. In contrast to the ε-VOPO4 prepared from the solvothermally synthesized H2VOPO4 either by heating or by electro-deintercalation of the protons, where very easy lithium

Figure 31. Cycling curves of ε-VOPO4 from (a) tetragonal and (b) monoclinic H2VOPO4 and (c) discharge curves as a function of current density of ε-VOPO4 from monoclinic H2VOPO4. Reprinted with permission from ref 159a. Copyright 2013 The Electrochemical Society.

characteristics of ε-VOPO4 formed from both the tetragonal and monoclinic forms of H2VOPO4.158 It can be seen that the capacity at the lower rates exceeds that of the olivines, but it is also apparent that there is a significant overpotential on charging. This might be associated with the preferred reaction path on recharge being one direct from Li2VOPO4 to ε-VOPO4 11433

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Figure 33. (top) Structure of Li2VOPO4 (two different orientations) and (bottom) showing changes in bond distances and angles with lithium content along the chain. Reprinted with permission from ref 171. Copyright 2014 Royal Society of Chemistry.

Figure 34. (Top) The almost open-circuit curves showing the redox potentials for Li1.5VOPO4, Li1.75VOPO4, and Li2VOPO4. (Middle) The cycling curves for the reaction LiVOPO4 to VOPO4. (Bottom) The steady increase of capacity and (inset) reduction of polarization on cycling. Reprinted with permission from ref 171. Copyright 2014 Royal Society of Chemistry.

insertion and removal were reported on the 4 V plateau, this LiVOPO4 showed very sluggish behavior. This is shown in the middle part of Figure 34. Only ∼50% of the lithium could be removed on the first charge. This steadily increases on continued cycling, and as shown in the lowest part of Figure 34, a constant value of ∼135 mAh/g is achieved after 25 cycles. The polarization also drops over these cycles. The authors attribute this to a breaking apart of the particles. This figure also shows an overcharge on each cycle, which may be due to electrolyte breakdown. Harrison and Manthiram174 synthesized LiVOPO4 by microwave solvothermal between 200 and 250 °C and formed products with a variety of sizes ranging from a few microns to 100 nm. The capacity attained reached up to 120− 134 mAh/g at the same C/20 rate on the first cycle even for some micron-size particles, so it is not clear that size alone is the cause of the low capacity exhibited in Figure 34. Closely related to LiVOPO4 is LiVPO4F, where the vanadyl oxygen is replaced by a fluoride ion. The initial work on this

compound was done by Barker et al.175 and in continuing papers showed a capacity of ∼120 mAh/g at a C/5 rate for 300 cycles.176 Mba et al. have made an in situ study of the lithium reactions with LiVOPO4.177 However, even with the higher voltage, this capacity is lower than that of LiFePO4, and LiVPO4F has not shown the same high rate capabilities. Thus, unless a second lithium can be incorporated, these tavorite-type compounds will see limited use as battery electrodes. Two redox potentials were first reported by Barker et al.178 for LiVPO4F, 4+ to 3+ and 3+ to 2+, but the split is even higher than that for the vanadyl phosphate, 4.2 and 1.8 V. They also described a symmetric cell with electrodes of LiVPO4F, which when charged became a cell as follows: VPO4F//Li2VPO4F. 11434

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The interested reader is referred to the recent review by Masquelier and Croguennec4 for further information on the tavorites. The beta phase has a structure close to that of ε-VOPO4. It is slightly more close-packed and is reported by Barker et al. to have a lower capacity at equivalent rates.162 Gaubicher et al.163 reported a capacity of 90 mAh/g (∼0.55 Li/VOPO4) for 80 cycles at the C/50 rate, but the capacity is only 40 mAh/g at C/ 4. Mixing it with RuO2 increased the capacity from 90 to 120 mAh/g.179 There are no reports of the intercalation of a second lithium into the structure. Dupré et al.180 investigated the electrochemical behavior of the hydrated α-VOPO4·nH2O phase and showed that it intercalates lithium by a standard single-phase reaction with a continuously decreasing potential. The capacity on cycling was improved after the insertion of sodium and ball-milling. The ω-VOPO4 phase is metastable, and no electrochemistry has been presented;165 the formula volume appears unrealistic (see Table 4), which might suggest the proposed structure is wrong. The sodium electrochemical behavior of some of the compounds has also been investigated.181 3.2. Other Multielectron Phosphates

Several phosphates beyond vanadium have been studied as potential hosts for lithium ions. None have yet been found, and most published effort has been placed on Fe, Mn, and Mo compounds. The iron and manganese pyrophosphates have been studied over the last 5 years by several groups. The electrochemistry and structure of Li2FeP2O7 was first reported by Zhou et al. at IMLB in 2010182 and independently studied by Yamada and co-workers.183 These compounds have the general formula Li2Fe1−yMnyP2O7, where there is a complete solid solution of the Fe and Mn. However, the Mn shows essentially no electrochemical activity at room temperature as shown in Figure 35, but one lithium reversibly cycles very well even though there appears to be a significant structural change. Some reactivity of the manganese compound is found at 40 °C and at a very low rate of C/50, showing that lithium diffusion is very sluggish. 184 Surprisingly, the sodium compound, Na2MnP2O7, shows significant electrochemical behavior, 90 mAh/g at room temperature, cycles well as shown in Figure 36,185 and is comparable to the iron analogue.184 The lithium cobalt compound, Li2CoP2O7, has also been studied186 and found to cycle, as shown in Figure 36c, but with a very large polarization on discharge and electrolyte decomposition on charge. Until electrolytes with a higher voltage window are found, interest in these pyrophosphates will wane. As can be observed from Figure 26, molybdenum phosphates should be good cathode candidates if all three oxidation states can be accessed. The electrochemical behavior of one molybdenum pyrophosphate, δ-(MoO2)P2O7, has been investigated by Wen et al;187 the originally synthesis procedure was followed.188 The electrochemical behavior is shown in Figure 37, and this indicates that although the capacity is relatively low the capacity retention in cycling is quite good. However, on every cycle there is an overcharge, which is surprising because the charging voltage is relatively low; this may indicate that this oxide has some undesired electrocatalytic activity. Figure 37c also indicates a falloff in capacity with increasing rate. Pushing the lithium content to four Li per formula unit gives a capacity of almost 250 mAh/g but turns the material amorphous. However, it remains very crystalline up to at least 1.2 Li; this composition is readily attained by reaction with LiI in

Figure 35. (top) Electrochemical behavior of Li2Mn1−yFeyP2O7 (y = 0, 0.2, 0.5, 0.8, and 1) at 0.1 mA/cm2182 and (bottom) capacity on cycling at various C rates of Li2FeyP2O7.182 Reprinted with permission from ref 182b. Copyright 2011 American Chemical Society.

acetonitrile, which has a redox potential of ∼3 V vs Li. The only other molybdenum phosphate, whose Li intercalation has been reported, is γ-(MoO2)2P2O7. Initially four Li ions can be intercalated, but it turns amorphous on charging to 4.2 V and loses capacity.189 The question is still open as to whether any molybdenum compound can store significant amounts of energy.

4. LAYERED OXIDES The layered oxides, LiMO2, have been the dominant cathode material for the last 20 years. Originally LiCoO2 was the dominant material,7 but over the past decade the mixed metal oxides, such as LiNi1/3Mn1/3Co1/3O2, have become more prevalent because of the cost of cobalt; the reader is referred to the 2004 review for more details of these materials.2 These so-called NMC materials are beginning to be used in electric vehicles, such as that produced by BMW. They have an inherently higher capacity than LiCoO2, a high rate capability, and lower cost;190 a complete understanding of the impact of 11435

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Figure 36. (a) Discharge curves as a function of rate (C/20, C/10, C/ 5, C/2, and C) and (b) capacity retention at C/5 for Na2MnP2O7.185 (c) Discharge curves as a function of rate for Li2CoP2O7; inset shows capacity for each cycle.186 Reprinted with permission from refs 185 and 186. Copyright 2013 and 2011, respectively, American Chemical Society.

metal composition on capacity and rate capability is still lacking. The high nickel content materials, LiNi0.80Co0.15Al0.05O2, known as NCA, are also finding application such as in the Panasonic cells used to power the Tesla electric vehicles. The aluminum in the structure prevents the complete oxidation of the nickel and, thus, reduces the oxygen chemical potential and increases the safety of the system. The status of these materials has not changed much since the review of 10 years ago,2 with the capacity still capped at ∼180 mAh/g of active material. These materials have intrinsically the highest theoretical energy density of around 1 kWh/kg and 3 kWh/L and probably represent the ultimate limit for intercalation reactions involving just one Li per redox center.

Figure 37. (a) Discharge−charge curves,187 (b) capacity retention on cycling,187 and (c) power capability of δ-(MoO2)P2O7.187 Reprinted with permission from ref 187. Copyright 2013 American Chemical Society.

4.1. Li- and Mn-Rich Layered Oxides

In the past decade there has been growing interest in the Lirich, Mn-rich oxides such as Li1.2[MnNiCo]0.8O2; they are also often described as Li2MnO3·LiMO2. They must be activated, as the Li2MnO3 component is electrochemically “inactive”, by an extended charge at ∼4.6 V. It is still not clear how much oxygen is evolved during the reaction or whether most of the oxygen is retained in the lattice and participates in the redox process. There has been an increased recognition that it is not only the transition metals but also the oxygen that participates in the redox process,191 as Rouxel had proposed for the transition metal disulfides.192 This material exhibits particularly high capacities at low rates and in particular at elevated temperatures, as shown in Figure 38.193 However, as can be seen from the figure, the voltage fades with cycling,194 probably indicating the formation of a spinel-like phase. Such a spinel phase has

Figure 38. Electrochemical profiles of 0.5Li 2 MnO 3 ·0.5LiNi0.44Mn0.31Co0.25O2 cathode (a) at room temperature and (b) at 55 °C. Reprinted with permission from ref 194. Copyright 2013 The Electrochemical Society.

been observed both in the bulk195 and on the surface as shown in Figure 39. As shown in Figure 39, a part of the transition metal ions enter the Li-layer in the bulk of the material, and a significant amount of the transition metals enter the Li-layers on the surface. This surface atomic rearrangement will 11436

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Figure 39. TEM of the surface region of a Li−Mn-rich layered oxide. (Left) Pristine material and (right) in the discharged state after being cycled 10 times between 2 and 4.8 V. Reprinted with permission from ref 195b. Copyright 2014 The Electrochemical Society.

significantly impede the diffusion of the lithium in and out of the structure. It is likely that this conversion is thermodynamically driven and, therefore, not readily prevented. There is a large ongoing applied and fundamental effort to overcome the challenges of these Li/Mn-rich materials, including a better understanding of the stoichiometric materials, LiMO2. This includes characterization efforts,196 studies of the arrangement including ordering of the ions, in particular manganese,197 in the lattice and how these change on cycling, and the impact of substitution in the lattice such as titanium for the expensive cobalt.198 4.2. Thermal Stability of Layered Oxides

A recurring concern with lithium batteries is their safety, and it is only magnified by each occurrence of a computer fire or the fire issues with the LiCoO2-based batteries on the Boeing 787. The layered oxides have long been known to tend to release oxygen on heating when fully charged, and aluminum is included in the NCA materials such as LiNi0.80Co0.15Al0.05O2 to prevent the full oxidation of nickel to 4+, i.e., the maximum oxygen chemical potential is limited. Their thermal stability on heating is not well understand, for example, why are some of the high manganese content materials less stable, when the Mn4+ ion is much more stable thermodynamically than Ni4+ or Co4+ in an oxide lattice? The answer might reside in the effective oxidation state of the oxygen anion. Noh et al.199 made a differential scanning calorimetry (DSC) study of NMC-type materials of formula Li1−δ[NixCoyMny]O2, showing (Figure 40) that the higher the nickel content the greater is the heat release and the lower is the temperature of the heat release, indicating the inherent instability of highnickel-containing layered oxides. As the residual lithium content of these phases varied from 0.34 for y = 0.33 to 0.21 for y = 0.075, all the peaks are probably on the high-temperature side of that expected for lithium-free materials. Choosing the optimal transition metal composition for these NMC materials presents a challenge. The higher the nickel content the greater is the storage capacity, but the capacity fades on cycling increases as does the thermal instability. The most commonly commercially used NMC material is the equimolar LiNi0.33Mn0.33Co0.33O2, which has the highest thermal stability and capacity retention but the lowest capacity. These trends are shown in Figure 40. Poor capacity retention and poor thermal stability appear to go hand-in-hand,

Figure 40. (Top) DSC heat evolution for a series of electrochemically delithiated NMC materials, Li1−δ[NixCoyMny]O2 where x = 1/3, 0.5, 0.6, 0.7, 0.8, and 0.85). (Bottom) Correlation of discharge capacity with capacity retention and thermal stability. Reprinted with permission from ref 199. Copyright 2013 Elsevier.

suggesting degradation of the material on cycling. This might be associated with reactions with the electrolyte or just with 11437

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oxygen loss. A study where the surface layer of high-nickel highly delithiated materials had a lower nickel content, 0.62− 0.67 Ni, than the core of the material particle of composition Li[Ni0.9Co0.05Mn0.05]O2 indeed showed that the coated materials had a greater thermal stability then the core material alone. This is indicated in Figure 41,199 and the heat evolved

Figure 41. DSC heat evolution for Li[Ni0.9Co0.05Mn0.05]O2 and it coated with lower nickel content phase, CSCG B − Ni = 0.62, and CSCG A − Ni = 0.67. Reprinted with permission from ref 199. Copyright 2013 Elsevier.

from the two coated examples CSCG B and CSCG A of 849 and 776 J/g, respectively, are significantly less than that, 1004 J/g, from the core material alone. These numbers are comparable to those shown in Figure 40. It has been generally assumed that adding manganese to these layered oxides will enhance the thermal stability, but the opposite is the case for the Mn-rich, Li-rich materials such as Li1.2Ni0.16Mn0.56Co0.08O2, which are significantly less thermally stable200 than NMC as shown in Figure 41. Substituting Al into the lattice improves the thermal stability of the lattice slightly, of similar amount to the coating shown in Figure 42a. The higher heat evolution shown in this work is probably related to the lower lithium contents of the material studied as the cells were charged to 4.6 V versus the 4.3 V of the previously discussed work. Although 5% aluminum substitution reduces the thermal instability, it reduces also the overall capacity as shown in Figure 42b, which shows very slow discharge rates and the rate capability as shown in Figure 42c. The last shows one of the other challenges faced by the Li−Mn-rich materials, that of low rates of reaction so that at high rates the traditional LiNi0.33Mn0.33Co0.33O2 has the best performance with much better thermal stability and capacity retention. There is little disagreement among reports on the thermal stability of the layered oxides. Little work has been reported on materials charged to 4.6 V, despite that need for the activation of the Li−Mn-rich oxides. There are several reports of Li-rich, but not Mn-rich (Mn < 0.5) materials. Kang et al. showed201 that the thermal stability of Li1.05(Mn1/3Co1/3Ni1/3)0.95O2 was greater than that of Li1.05(Mn4/9Co1/9Ni4/9)0.95O2 when charged to 4.4 and 4.3 V, respectively. Croguennec et al. showed202 similar results for Li1+x(Mn0.4Ni0.4Co0.2)1−xO2 versus NCA, when charged to 4.5 V and x was LiNi 2/3 Mn 1/3 O 2 > LiCoO 2 > LiNi0.5Mn1.5O4 > Li[Li0.11Ni0.33Mn0.56]O2 > Li[Li0.2Ni0.2Mn0.6]O2. The Li-rich are the least stable and all three materials where the Mn content was >50% of the transition metal content performed the poorest. These results all tend to suggest surprisingly that it is the Mn-rich component of the Li−Mnrich materials that is leading to the thermal instability. This is contrary to the original reports where it was thought that increasing the manganese content would improve the thermal stability.205 11438

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these that is commercial is Li−S, where Sion is producing Li−S cells of limited cycling capability with storage levels of 390 Wh/ kg and 390 Wh/L.207 These two numbers show the strength and weakness of nonintercalation electrodes, namely, the potential to perhaps double the gravimetric storage capability of batteries but at the expense of volume storage capability, which may be not much better than that of comparable intercalation batteries. Li−S cells are reviewed in this issue by Manthiram et al. Li−O2 cells, which have attractive gravimetric capacities and are reviewed by Luntz and McCloskey in this issue, may have insurmountable technical challenges, whose cost to overcome may not be warranted by the calculated incremental improvements possible. In conclusion, it is probably fair to say that the rate of intercalation reactions is quite sufficient for most storage and power applications, being able to be discharged or charged in ≪1 h. The ultimate limit to the storage capabilities of intercalation reactions will be around 0.5 kWh/kg and 1.5 Wh/L in full cell configurations. Alternate reaction schemes will be needed to achieve higher gravimetric storage levels, and it is unlikely that any other battery will exceed the volumetric storage capability of intercalation-based cells at the full battery level.

It should obviously be noted that the cathode is not the only contributor to safety issues in lithium batteries, but a thermally stable material is clearly superior, particularly one that does not generate oxygen when charged within or beyond normal limits. The electrolyte is a key safety issue, providing the fuel for many battery incidences. However, the anode and in particular today’s carbon-based anodes are a key source of fuel and pressure buildup if combusted. Table 5 shows the energy of Table 5. Free Energy of Combustion of Anode Materials material

combustion product

energy, kJ/Li

C Si Sn2Fe

CO2 SiO2 SnO2 + Fe2O3

2366 194 160

combustion of carbon compared with the potential nextgeneration anodes, silicon and tin; the data is presented per Li stored. It is clear that carbon poses a much higher safety risk.

5. CONCLUSIONS AND WHAT DOES THE FUTURE HOLD This review described predominantly two model intercalation compounds, LiFePO4 and VOPO4, to explore the mechanisms of battery electrode reactions. Comparing their behavior with that of the layered oxides gives a view of the limitations of Liion batteries based on intercalation reactions. Intercalationbased reactions provide the opportunity for the highest rates in solid electrodes because there is no need to build structures on discharge and charge. However, their capacity is limited by the fact that a host material is required to hold the lithium, as well as that the host material has mass and volume. The ultimate storage limit for active materials alone is ∼1 kWh/kg and 3 kW/L. Practical cells should be able to attain 50% of those values with optimized engineering design. Attainment of these goals will require materials with faster ion transport so that thicker electrodes can be used; that will allow the elimination of at least 50% of the current collector and separator area, and thus volume, weight, and cost. One critical component not discussed is the graphitic carbon anode that reacts, forming the intercalation compound, C6Li. This anode takes up half the cell volume and weighs 84 g/mol of Li stored. Elimination of the carbon is needed to attain the ultimate storage limit for the intercalation cathode. The ideal is to use lithium metal, thus avoiding any dead weight. The impacts on the volumetric energy density of intercalation batteries of using lithium have recently been highlighted by Gallagher et al.206 Several materials approaches have been used as a step toward moving toward higher capacity anodes, principally involving predominantly aluminum (as used in the first Li-ion batteries built by Exxon), tin, and silicon. These are reviewed in this volume by Obravac and Chevrier (DOI 10.1021/cr500207g). Another approach to reaching the ultimate limit of intercalation batteries is to replace Li by Mg, which is expected to have a higher volumetric capacity, Ah/kg, than corresponding 2 Li systems, but perhaps at the expense of some of the operating voltage. Mg systems are reviewed in this issue by Muldoon et al (DOI 10.1021/cr500049y). There has been much recent work on battery systems exploring alternatives to intercalation electrodes. This special issue covers two of these approaches, Li−O2 (DOI 10.1021/ cr500054y) and Li−S (DOI 10.1021/cr500062v); a later tissue will cover conversion cathodes such as FeOF. The only one of

AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest. Biography

M. Stanley Whittingham was born in Nottingham, England, and received his B.A. and D.Phil. degrees in Chemistry from Oxford University working with Peter Dickens. In 1968 he went to Professor Robert A. Huggins’s research group in the Materials Science Department at Stanford University as a Postdoctoral Research Associate to study fast-ion transport in solids. In 1972 he joined Exxon Research and Engineering Company to initiate a program in alternative energy production and storage. After 16 years in industry he joined the Binghamton campus of the State University of New York to initiate an academic program in Materials Chemistry. Presently he is also Distinguished Professor of Chemistry and Materials and Director of the Materials Science Program and Institute for Materials Research. He was awarded the Young Author Award of the Electrochemical Society in 1971, a JSPS Fellowship in the Physics Department of the University of Tokyo in 1993, and the Battery Research Award of the Electrochemical Society in 2002 and was elected a Fellow of the Electrochemical Society in 2004 and of the Materials Research Society 11439

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(7) Mitzushima, K.; Jones, P. C.; Wiseman, P. J.; Goodenough, J. B. Mater. Res. Bull. 1980, 15, 783. (8) Padhi, A. K.; Nanjundaswamy, K. S.; Goodenough, J. B. J. Electrochem. Soc. 1997, 144, 1188. (9) (a) Ravet, N.; Goodenough, J. B.; Besner, S.; Simoneau, M.; Hovington, P.; Armand, M. Electrochem. Soc. Abstr. 1999, 99-2, 127. (b) Ravet, N.; Besner, S.; Simoneau, M.; Vallée, A.; Armand, M.; Magnan, J.-F. European Patent 1049182A2, 2000. (c) Ravet, N.; Chouinard, Y.; Magnan, J. F.; Besner, S.; Gauthier, M.; Armand, M. J. Power Sources 2001, 97−98, 503. (10) Wang, Q.; Zakeeruddin, S. M.; Wang, D.; Exnar, I.; Graetzel, M. Angew. Chem. 2006, 45, 8197. (11) Chung, S.-Y.; Bloking, J. T.; Chiang, Y.-M. Nat. Mater. 2002, 1, 123. (12) Chen, J.; Vacchio, M. J.; Wang, S.; Chernova, N.; Zavalij, P. Y.; Whittingham, M. S. Solid State Ionics 2008, 178, 1676. (13) Kamauchi, M.; Soejima, H.; Kubota, S.; Sasaki, K. U.S. Patent 5,538,814, July 23, 1996. (14) Yakubovich, O. V.; Bairakov, V. V.; Simonov, M. A. Dokl. Akad. Nauk SSSR 1989, 307, 1119. (15) Libowitzky, E.; Beran, A.; Wieczorek, A. K.; Wirth, R. Mineral. Pet. 2012, 105, 3. (16) Belov, N. V.; Matvienko, E. N.; Ivashchenko, A. N.; Yakubovich, O. V.; Mel’nikov, O. K.; Simonov, M. A. Mineral. Zh. 1981, 3, 56. (17) (a) Shoji, H.; Toshikazu, K.; Toshimi, F.; Makoto, F. Japanese Patent Publication 2002-151082, 2002. (b) Nuspl, G.; Wimmer, L.; Eisgruber, M. Canadian Patent Application CA2537278, 2005. (c) Nuspi, G.; Vogler, C.; Eisgruber, M.; Wimmer, L.; Schall, N.; Fietzek, C.; Weydanz, W. 12th Int. Meet. Lithium Batteries 2004, 12, 293. (d) Schall, N.; Nuspl, G.; Christian, V.; Wimmer, L.; Eisgruber, M. European Patent WO2006105848, 2006. (18) Yang, S.; Zavalij, P. Y.; Whittingham, M. S. Electrochem. Commun. 2001, 3, 505. (19) Yang, S.; Song, Y.; Zavalij, P. Y.; Whittingham, M. S. Electrochem. Commun. 2002, 4, 239. (20) Chen, J.; Whittingham, M. S. Electrochem. Commun. 2006, 8, 855. (21) Axmann, P.; Stinner, C.; Wohlfahrt-Mehrens, M.; Mauger, A.; Gendron, F.; Julien, C. M. Chem. Mater. 2009, 21, 1636. (22) Yakubovich, O. V.; Simonov, M. A.; Belov, N. V. Soviet Phys.Dokl. 1977, 22, 347. (23) Streltsov, V. A.; Belokoneva, E. L.; Tsirelson, V. G.; Hansen, N. K. Acta Crystallogr. 1993, B49, 147. (24) Goñi, A.; Bonagamba, T. J.; M. A. Silva, H. P.; Rojo, T.; Barberis, G. E. J. Appl. Phys. 1998, 84, 416. (25) Goñi, A.; Lezama, L.; Arriortua, M. I.; Barberis, G. E.; Rojo, T. J. Mater. Chem. 2000, 10, 423. (26) Ellis, B.; Kan, W. H.; Makahnouk, W. R. M.; Nazar, L. F. J. Mater. Chem. 2007, 17, 3248. (27) Chen, J.; Bai, J.; Chen, H.; Graetz, J. J. Phys. Chem. Lett. 2011, 2, 1874. (28) Chen, J. J.; Graetz, J. ACS Appl. Mater. Interfaces 2011, 3, 1380. (29) Jensen, K.; Christensen, M.; Tyrsted, C.; Iversen, B. B. J. Appl. Crystallogr. 2011, 44, 287. (30) Qin, X.; Wang, X.; Xiang, H.; Xie, J.; Li, J.; Zhou, Y. J. Phys. Chem. C 2010, 114, 16806. (31) Jensen, K. M. Ø.; Christensen, M.; Gunnlaugsson, H. P.; Lock, N.; Bøjesen, E. D.; Proffen, T.; Iversen, B. B. Chem. Mater. 2013, 25, 2282. (32) Murugan, A. V.; Muraliganth, T.; Manthiram, A. Electrochem. Commun. 2008, 10, 903. (33) Murugan, A. V.; Muraliganth, T.; Manthiram, A. J. Electrochem. Soc. 2009, 156, A79. (34) Delacourt, C.; Poizot, P.; Levasseur, S.; Masquelier, C. Electrochem. Solid-State Lett. 2006, 9, A352. (35) Recham, N.; Armand, M.; Laffont, L.; Tarasconz, J.-M. Electrochem. Solid-State Lett. 2009, 12, A39. (36) Chen, J. Recent Pat. Nanotechnol. 2013, 7, 2.

in 2013. He received from IBA the Yeager Award for Lifetime Contributions to Lithium Battery Materials Research in 2012 and in 2010 the ACS NERM Award for Contributions to Chemistry. He was Principal Editor of the journal Solid State Ionics for 20 years. His recent work focuses on the synthesis and characterization of novel materials for battery electrodes and he leads the DOE NECCES EFRC.

ACKNOWLEDGMENTS This work was supported by the U.S. Department of Energy, Office of Science, under Award numbers DE-SC0001294 and DE-SC0012583, through the NorthEast Center for Chemical Energy Storage (NECCES), an Energy Frontier Research Center, with collaborators on intercalation materials at Argonne National Laboratory (Karena Chapman), Binghamton University (Natasha Chernova, Frederik Omenya), Brookhaven National Laboratory (Jason Graetz, Xiao-Qing Yang), Cambridge University (Clare Grey), MIT (Gerbrand Ceder), Stony Brook University (Peter Khalifah), and U.C. San Diego (Shirley Meng) whose research is included in this review. I thank the above for many fruitful discussions. In addition I thank Ruibo Zhang for literature evaluation of the VOPO4 structures, and Natasha Chernova for the contents figure. ABBREVIATIONS anode cathode C rate

AND SPECIALIZED TERMS electropositive electrode electronegative electrode measure of the time for cell discharge in reciprocal hours, time of discharge = 1/C (h) Ellingham diagram diagram describing the stability of, in particular, oxides with temperature and oxygen partial pressure HEV hybrid electric vehicle LiPF6 lithium hexafluorophosphate, LiPF6 lithiophyllite mineral name for lithium-free Fe1−yMnyPO4 NASICON sodium superionic conductor; also used to describe the structural type NCA lithium nickel cobalt aluminum oxide NMC lithium nickel manganese cobalt oxide Ragone plot plot of the electrochemical cell capacity as a function of the magnitude of the discharge or charge current Rietveld analysis specialized computer technique to resolve the structure of compounds from powder diffraction data simferite naturally occurring mineral with the olivine structure containing iron, manganese, and magnesium tryphilite mineral name for the lithium-containing LiFe1−yMnyPO4 Vegard’s law for a solid solution, the lattice parameter varies linearly with composition REFERENCES (1) Whittingham, M. S. Proc. IEEE 2012, 100, 1518. (2) Whittingham, M. S. Chem. Rev. 2004, 104, 4271. (3) Fletcher, S. Bottled Lightning: Superbatteries, Electric Cars, and the New Lithium Economy; Hill and Wang: New York, 2011. (4) Masquelier, C.; Croguennec, L. Chem. Rev. 2013, 113, 6552. (5) Winter, M.; Brodd, R. J. Chem. Rev. 2004, 104, 4245. (6) Whittingham, M. S. Science 1976, 192, 1126. 11440

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Review

(66) Delacourt, C.; Poizot, P.; Tarascon, J. M.; Masquelier, C. Nat. Mater. 2005, 4, 254. (67) Bramnik, N. N.; Nikolowski, K.; Baehtz, C.; Bramnik, K. G.; Ehrenberg, H. Chem. Mater. 2007, 19, 908. (68) Zhu, Y.; Wang, J. W.; Liu, Y.; Liu, X.; Kushima, A.; Liu, Y.; Xu, Y.; Mao, S. X.; Li, J.; Wang, C.; Huang, J. Y. Adv. Mater. 2014, 24, 312. (69) Allen, J.; Jow, T.; Wolfenstine, J. Chem. Mater. 2007, 19, 2108. (70) Orikasa, Y.; Maeda, T.; Koyama, Y.; Murayama, H.; Fukuda, K.; Tanida, H.; Arai, H.; Matsubara, E.; Uchimoto, Y.; Ogumi, Z. Chem. Mater. 2013, 25, 1032. (71) Liu, H.; Strobridge, F. C.; Borkiewicz, O. J.; Wiaderek, K. M.; Chapman, K. W.; Chupas, P. J.; Grey, C. P. Science 2014, 344, 1252817. (72) Zhang, X.; Hulzen, M. v.; Singh, D. P.; Brownrigg, A.; Wright, J. P.; Dijk, N. H. v.; Wagemaker, M. Nano Lett. 2014, 14, 2279. (73) Borkiewicz, O. J.; Shyam, B.; Wiaderek, K. M.; Kurtz, C.; Chupas, P. J.; Chapman, K. W. J. Appl. Crystallogr. 2012, 45, 1261. (74) Bai, P.; Cogswell, D. A.; Bazant, M. Z. Nano Lett. 2011, 11, 4890. (75) Srinivasan, V.; Newman, J. Electrochem. Solid-State Lett. 2006, 9, A110. (76) Delmas, C.; Maccario, M.; Croguennec, L.; Cras, F. L.; Weill, F. Nat. Mater. 2008, 7, 665. (77) Brunetti, G.; Robert, D.; Bayle-Guillemaud, P.; Rouviere, J. L.; Rauch, E. F.; Martin, J. F.; Colin, J. F.; Bertin, F.; Cayron, C. Chem. Mater. 2011, 23, 4515. (78) Wagemaker, M.; Singh, D. P.; Borghols, W. J. H.; Lafont, U.; Haverkate, L. A.; Peterson, V. K.; Mulder, F. M. J. Am. Chem. Soc. 2011, 133, 10222. (79) Dedryvère, R.; Maccario, M.; Croguennec, L.; Cras, F. L.; Delmas, C.; Gonbeau, D. Chem. Mater. 2008, 20, 7164. (80) Wang, F.; Robert, R.; Chernova, N. A.; Pereira, N.; Omenya, F.; Badway, F.; Hua, X.; Ruotolo, M.; Zhang, R.; Wu, L.; Volkov, V.; Su, D.; Key, B.; Whittingham, M. S.; Grey, C. P.; Amatucci, G. G.; Zhu, Y.; Graetz, J. J. Am. Chem. Soc. 2011, 133, 18828. (81) Weichert, K.; Sigle, W.; Aken, P. A. v.; Jamnik, J.; Zhu, C.; Amin, R.; Acartürk, T.; Starke, U.; Maier, J. J. Am. Chem. Soc. 2012, 134, 2988. (82) Chueh, W. C.; Gabaly, F. E.; Sugar, J. D.; Bartelt, N. C.; McDaniel, A. H.; Fenton, K. R.; Zavadil, K. R.; Tyliszczak, T.; Lai, W.; McCarty, K. F. Nano Lett. 2013, 13, 866. (83) Malik, R.; Zhou, F.; Ceder, G. Nat. Mater. 2011, 10, 587. (84) Dreyer, W.; Jamnik, J.; Guhlke, C.; Huth, R.; Moškon, J.; ̌ Gaberšcek, M. Nat. Mater. 2010, 10, 448. (85) Meethong, N.; Kao, Y.-H.; Tang, M.; Huang, H.-Y.; Carter, W. C.; Chiang, Y.-M. Chem. Mater. 2008, 20, 6189. (86) Yoo, S.; Kang, B. Nanotechnology 2013, 24, 424012. (87) Sugar, J. D.; Gabaly, F. E.; Chueh, W.; Fenton, K. R.; Tyliszczak, T.; G, P.; Kotula; Bartelt, N. C. J. Power Sources 2014, 246, 512. (88) Love, C. T.; Korovina, A.; Patridge, C. J.; Swider-Lyons, K. E.; Twigg, M. E.; Ramakera, D. E. J. Electrochem. Soc. 2013, 160, A3153− A3161. (89) Malik, R.; Abdellahi, A.; Ceder, G. J. Electrochem. Soc. 2013, 160, A3179. (90) Xiao, J.; Yu, X.; Zheng, J.; Zhou, Y.; Gao, F.; Chen, X.; Bai, J.; Yang, X.-Q.; Zhang, J.-G. J. Power Sources 2013, 242, 736. (91) Whittingham, M. S. Mater. Res. Bull. 1978, 13, 959. (92) Ellis, B.; Herle, P. S.; Rho, Y.-H.; Nazar, L. F.; Dunlap, R.; Perry, L. K.; Ryan, D. H. Faraday Discuss. 2007, 134, 119. (93) Omenya, F.; Chernova, N. A.; Upreti, S.; Zavalij, P. Y.; Nam, K.W.; Yang, X.-Q.; Whittingham, M. S. Chem. Mater. 2011, 23, 4733. (94) Omenya, F.; Chernova, N. A.; Wang, Q.; Zhang, R.; Whittingham, M. S. Chem. Mater. 2013, 25, 2691. (95) Meethong, N.; Kao, Y.-H.; Speakman, S. A.; Chiang, Y.-M. Adv. Funct. Mater. 2009, 19, 1. (96) Delacourt, C.; C. Wurm, L. L.; Leriche, J.-B.; Masquelier, C. Solid State Ionics 2006, 177, 333. (97) Wagemaker, M.; Ellis, B. L.; Lützenkirchen-Hecht, D.; Mulder, F. M.; Nazar, L. F. Chem. Mater. 2008, 20, 6313.

(37) Chernova, N.; Nolis, G. M.; Omenya, F. O.; Zhou, H.; Li, Z.; Whittingham, M. S. J. Mater. Chem. 2011, 21, 9865. (38) Santoro, R. P.; Newnham, R. E. Acta Crystallogr. 1967, 22, 344. (39) Rousse, G.; Rodriguez-Carvajal, J.; Patoux, S.; Masquelier, C. Chem. Mater. 2003, 15, 4082. (40) Santoro, R. P.; Segal, D. J.; Newnham, R. E. J. Phys. Chem. Solids 1966, 27, 1192. (41) Rissouli, K.; Benkhouja, K.; Bettach, M.; Sadel, A.; Zahir, M.; Derrory, A.; Drillon, M. Ann. Chim. Sci. Mater. 1998, 23, 85. (42) (a) Barker, J.; Saidi, M. Y.; Swoyer, J. L. Electrochem. Solid-State Lett. 2003, 6, A53. (b) Barker, J.; Saidi, M. Y.; Swoyer, J. U.S. Patent 6,528,033, March 4, 2003. (c) Barker, J.; Saidi, M. Y. U.S. Patent 7,001,690, Feb 21, 2006. (d) Barker, J.; Saidi, M. Y.; Swoyer, J. L. J. Electrochem. Soc. 2003, 150, A684. (43) Zheng, J.-C.; Li, X.-H.; Wang, Z.-X.; Guo, H.-J.; Zhou, S.-Y. J. Power Sources 2008, 184, 574. (44) (a) Gauthier, M.; Michot, C.; Ravet, N.; Duchesneau, M.; Dufour, J.; Liang, G.; Wontcheu, J.; Gauthier, L.; MacNeilb, D. D. J. Electrochem. Soc. 2010, 157, A453. (b) MacNeil, D. D.; Devigne, L.; Michot, C.; Liang, I. R.; Gauthier, M. J. Electrochem. Soc. 2010, 157, A463. (45) Kim, J.-K.; Scheers, J.; Hwang, G.-C.; Zhao, X.; Kang, S.; Johansson, P.; Ahna, J.-H.; Jacobsson, P. Scr. Mater. 2013, 69, 716. (46) Dodd, J. L.; Yazami, R.; Fultz, B. Electrochem. Solid-State Lett. 2006, 9, A151. (47) Zhu, C.; Gu, L.; Suo, L.; Popovic, J.; Li, H.; Ikuhara, Y.; Maier, J. Adv. Funct. Mater. 2014, 3, 312. (48) (a) Yamada, A.; Koizumi, H.; Nishimura, S. I.; Sonoyama, N.; Kanno, R.; Yonemura, T.; Kobayashi, O. Nat. Mater. 2006, 5, 357. (b) Yamada, A.; Koizumi, H.; Sonoyama, N.; Kanno, R. Electrochem. Solid-State Lett. 2005, 8, A409. (c) Yamada, A.; Kudo, Y.; Liu, K.-Y. J. Electrochem. Soc. 2001, 148, A1153. (49) Meethong, N.; Huang, H.-Y. S.; Carter, W. C.; Chiang, Y.-M. Electrochem. Solid-State Lett. 2007, 10, A134. (50) Kim, D.-H.; Kim, J. Electrochem. Solid-State Lett. 2006, 9, A439. (51) Ong, S. P.; Wang, L.; Kang, B.; Ceder, G. Chem. Mater. 2008, 20, 1798. (52) Janssen, Y.; Santhanagopalan, D.; Qian, D.; Chi, M.; Wang, X.; Hoffmann, C.; Meng, Y. S.; Khalifah, P. Chem. Mater. 2013, 25, 4574. (53) Chianelli, R. R.; Scanlon, J. C.; Rao, B. M. L. J. Solid State Chem. 1979, 29, 323. (54) Islam, M. S.; Driscoll, D. J.; Fisher, C. A. J.; Slater, P. R. Chem. Mater. 2005, 17, 5085. (55) Nishimura, S.-I.; Kobayashi, G.; Ohoyama, K.; Kanno, R.; Yashima, M.; Yamada, A. Nat. Mater. 2008, 7, 707. (56) Chen, G.; Song, X.; Richardson, T. J. Electrochem. Solid-State Lett. 2006, 9, A295. (57) Laffont, L.; Delacourt, C.; Gibot, P.; Wu, M. Y.; Kooyman, P.; Masquelier, C.; Tarascon, J. M. Chem. Mater. 2006, 18, 5520. (58) Liu, X.; Wang, D.; Liu, G.; Srinivasan, V.; Liu, Z.; Hussain, Z.; Yang, W. Nat. Commun. 2013, 4, 2568. (59) Chen, G.; Song, X.; Richardson, T. J. J. Electrochem. Soc. 2007, 154, A627. (60) Gu, L.; Zhu, C.; Li, H.; Yu, Y.; Li, C.; Tsukimoto, S.; Maier, J.; Ikuhara, Y. J. Am. Chem. Soc. 2011, 133, 4661. (61) Robert, D.; Douillard, T.; Boulineau, A.; Brunetti, G.; Nowakowski, P.; Venet, D.; Bayle-Guillemaud, P.; Cayron, C. ACS Nano 2013, 12, 10887. (62) Omenya, F. Olivine substitution and reaction mechanism. Ph.D. Dissertation, Binghamton University, Binghamton, NY, 2013. (63) Omenya, F.; Chernova, N. A.; Zhang, R.; Fang, J.; Huang, Y.; Cohen, F.; Dobrzynski, N.; Senanayake, S.; Xu, W.; Whittingham, M. S. Chem. Mater. 2013, 25, 85. (64) Park, K.-Y.; Hong, J.; Kim, J.; Park, Y.-U.; Kim, H.; Seo, D.-H.; Kim, S.-W.; Choi, J.-W.; Kanga, K. J. Electrochem. Soc. 2013, 160, A444. (65) Gibot, P.; Casas-Cabanas, M.; Laffont, L.; Levasseur, S.; Carlach, P.; Hamelet, S.; Tarascon, J.-M.; Masquelier, C. Nat. Mater. 2008, 7, 741. 11441

dx.doi.org/10.1021/cr5003003 | Chem. Rev. 2014, 114, 11414−11443

Chemical Reviews

Review

(98) Harrison, K. L.; Bridges, C. A.; Paranthaman, M. P.; Segre, C. U.; Katsoudas, J.; Maroni, V. A.; Idrobo, J. C.; Goodenough, J. B.; Manthiram, A. Chem. Mater. 2013, 25, 768. (99) Islam, M. S.; Driscoll, D. J.; Fisher, C. A. J.; Slater, P. R. Chem. Mater. 2005, 17, 5085. (100) Herle, P. S.; Ellis, B.; Coombs, N.; Nazar, L. F. Nat. Mater. 2004, 3, 147. (101) Shi, S.; Ouyang, C.; Wang, D.-S.; Wang, Z.; Chen, L.; Huang, X. Phys. Rev. 2003, B68, 195108. (102) (a) Ravet, N.; Abouimrane, A.; Armand, M. Nat. Mater. 2003, 2, 702. (b) Chung, S.-Y.; Bloking, J. T.; Chiang, Y.-M. Nat. Mater. 2003, 2, 702. (103) Liu, H.; Xie, J.; Wang, K. Solid State Ionics 2008, 179, 1768. (104) Zhou, X.; Wang, F.; Zhu, Y.; Liu, Z. J. Mater. Chem. 2011, 21, 3353. (105) Amin, R.; Balaya, P.; Maier, J. Electrochem. Solid-State Lett. 2007, 10, A13. (106) Hong, J.; Wang, C. S.; Chen, X.; Upreti, S.; Whittingham, M. S. Electrochem. Solid-State Lett. 2009, 12, A33. (107) Hua, N.; Wang, C.; Kang, X.; Wumair, T.; Han, Y. J. Alloys Compd. 2010, 503, 204. (108) Ma, J.; Li, B.; Du, H.; Xu, C.; Kanga, F. J. Electrochem. Soc. 2011, 158, A26. (109) Sun, C. S.; Zhou, Z.; Xu, Z. G.; Wang, D. G.; Wei, J. P.; Bian, X. K.; Yan, J. J. Power Sources 2009, 193, 841. (110) Zhang, L.-L.; Liang, G.; Ignatov, A.; Croft, M. C.; Xiong, X.-Q.; Hung, I.-M.; Huang, Y.-H.; Hu, X.-L.; Zhang, W.-X.; Peng, Y.-L. J. Phys. Chem. C 2011, 115, 13520. (111) Jin, Y.; Yang, C. P.; Rui, X. H.; Cheng, T.; Chen, C. H. J. Power Sources 2011, 196, 5623. (112) LFP-NCO Cathode Material. http://www.aleees.com (accessed August 24, 2014). (113) Zhao, T.; Xu, W.; Ye, Q.; Cheng, J.; Zhao, H. F.; Wu, Z. Y.; Xia, D. G.; Chu, W. S. J. Synchrotron Radiat. 2010, 17, 584. (114) Harrison, K. L.; Manthiram, A. Inorg. Chem. 2011, 50, 3613. (115) Hong, J.; Wang, X.-L.; Wang, Q.; Omenya, F.; Chernova, N. A.; Whittingham, M. S.; Graetz, J. J. Phys. Chem. C 2012, 116, 20787. (116) Meethong, N.; Kao, Y.-H.; Carter, W. C.; Chiang, Y.-M. Chem. Mater. 2010, 22, 1088. (117) Hua, Y.; Yao, J.; Zhao, Z.; Zhu, M.; Li, Y.; Jin, H.; Zhao, H.; Wang, J. Mater. Chem. Phys. 2013, 141, 835. (118) Zeng, G.; Caputo, R.; Carriazo, D.; Luo, L.; Niederberger, M. Chem. Mater. 2013, 25, 3399. (119) Dodd, J. Thermodynamics of Iron Phosphates? Ph.D. Dissertation, California Institute of Technology, Pasadena, CA, 2007. (120) Kim, J.; Park, K.-Y.; Park, I.; Yoo, J. K.; Hong, J.; Kang, K. J. Mater. Chem. 2012, 22, 11964. (121) Yang, S.; Song, Y.; Ngala, K.; Zavalij, P. Y.; Whittingham, M. S. J. Power Sources 2003, 119, 239. (122) Shu, J.; Shui, M.; Xu, D.; Wang, D.; Ren, Y.; Gao, S. J. Solid State Electrochem. 2012, 16, 819. (123) He, H.; Liu, Y.; Liu, Q.; Li, Z.; Xu, F.; Dun, C.; Ren, Y.; Wang, M.-x.; Xie, J. J. Electrochem. Soc. 2013, 160, A793. (124) Koltypin, M.; Aurbach, D.; Nazar, L.; Ellis, B. Electrochem. Solid-State Lett. 2007, 10, A40. (125) Guéguen, A.; Castro, L.; Dedryvère, R.; Dumont, E.; Bréger, J.; Tessier, C.; Gonbeau, D. J. Electrochem. Soc. 2013, 160, A387. (126) Hamelet, S. e.; Gibot, P.; Casas-Cabanas, M.; Bonnin, D.; Grey, C. P.; Cabana, J.; Leriche, J.-B.; Rodriguez-Carvajal, J.; Courty, M.; Levasseur, S. E.; Thournout, P. C. E. V.; Tarascon, J.-M.; Masquelier, C. J. Mater. Chem. 2009, 19, 3979. (127) Amisse, R.; Hamelet, S. e.; Hanzel, D.; Courty, M.; Dominko, R.; Masquelier, C. J. Electrochem. Soc. 2013, 160, A1446. (128) Chen, G.; Richardson, T. J. J. Power Sources 2010, 195, 1221. (129) Ong, S. P.; Jain, A.; Hautier, G.; Kang, B.; Ceder, G. Electrochem. Commun. 2010, 12, 427. (130) Nolis, G.; Omenya, F.; Zhang, R.; Fang, B.; Upreti, S.; Chernova, N.; Wang, F.; Graetz, J.; Hu, Y.-Y.; Grey, C.; Whittingham, M. S. J. Mater. Chem. 2012, 22, 20482.

(131) Huang, Y.; Fang, J.; Omenya, F.; O’Shea, M.; Chernova, N. A.; Zhang, R.; Wang, Q.; Quackenbush, N. F.; Piper, L. F. J.; Scanlon, D. O.; Whittingham, M. S. J. Mater. Chem. A 2014, 2, 12827. (132) Delacourt, C.; Laffont, L.; Bouchet, R.; Wurm, C.; Leriche, J. B.; Morcrette, M.; Tarascon, J. M.; Masquelier, C. J. Electrochem. Soc. 2005, 152, A913. (133) Yamada, A.; Chung, S.-C. J. Electrochem. Soc. 2001, 148, A960. (134) Clément, R. l. J.; Pell, A. J.; Middlemiss, D. S.; Strobridge, F. C.; Miller, J. K.; Whittingham, M. S.; Emsley, L.; Grey, C. P.; Pintacuda, G. J. Am. Chem. Soc. 2012, 134, 17178. (135) Piper, L. F. J.; Quackenbush, N. F.; Sallis, S.; Scanlon, D. O.; Watson, G. W.; Nam, K.-W.; Yang, X.-Q.; Smith, K. E.; Omenya, F.; Chernova, N. A.; Whittingham, M. S. J. Phys. Chem. C 2013, 117, 10383. (136) Kwon, N.-H.; Drezen, T.; Exnar, I.; Teerlinck, I.; Isono, M.; Graetzel, M. Electrochem. Solid-State Lett. 2006, 9, A277. (137) Drezen, T.; Kwon, N.-H.; Bowenb, P.; Teerlinck, I.; Isono, M.; Exnar, I. J. Power Sources 2007, 174, 949. (138) Delacourt, C.; Poizot, P.; Morcrette, M.; Tarascon, J.-M.; Masquelier, C. Chem. Mater. 2004, 16, 93. (139) Yonemura, M.; Yamada, A.; Takei, Y.; Sonoyama, N.; Kanno, R. J. Electrochem. Soc. 2004, 151, A1352. (140) Wang, D.; Ouyang, C.; Drézen, T.; Exnar, I.; Kay, A.; Kwon, N.-H.; Gouerec, P.; Miners, J. H.; Wang, M.; Grätzela, M. J. Electrochem. Soc. 2010, 157 (2), A225. (141) Ravnsbæk, D. B.; Xiang, K.; Xing, W.; Borkiewicz, O. J.; Wiaderek, K. M.; Gionet, P.; Chapman, K. W.; Chupas, P. J.; Chiang, Y.-M. Nano Lett. 2014, 14, 1484. (142) Malik, R.; Zhou, F.; Ceder, G. Phys. Rev. B 2009, 79, 214201. (143) Amine, K.; Yasuda, H.; Yamachib, M. Electrochem. Solid-State Lett. 2000, 3, 178. (144) Okada, S.; Sawa, S.; Egashira, M.; Yamaki, J.-i.; Tabuchi, M.; Kageyama, H.; Konishi, T.; Yoshino, A. J. Power Sources 2001, 97−98, 430. (145) Nakayama, M.; Goto, S.; Uchimoto, Y.; Wakihara, M.; Kitajima, Y. Chem. Mater. 2004, 16 (18), 3399. (146) Shui, J. L.; Yu, Y.; Yang, X. F.; Chen, C. H. Electrochem. Commun. 2006, 8, 1087. (147) Eftekhari, A. J. Electrochem. Soc. 2004, 151, A1456. (148) Li, H. H.; Jin, J.; Wei, J. P.; Zhou, Z.; Yan, J. Electrochem. Commun. 2008, 11, 95. (149) Han, D.-W.; Kang, Y.-M.; Yin, R.-Z.; Song, M.-S.; Kwon, H.-S. Electrochem. Commun. 2008, 11, 137. (150) Truong, Q. D.; Devaraju, M. K.; Tomai, T.; Honma, I. ACS Appl. Mater. Interfaces 2013, 5, 9926. (151) Ehrenberg, H.; Bramnik, N. N.; Senyshyn, A.; Fuess, H. Solid State Sci. 2008, 10, 1. (152) Ni, J.; Liu, W.; Liua, J.; Gao, L.; Chen, J. Electrochem. Commun. 2013, 35, 1. (153) Lee, H.; Kim, M. G.; Cho, J. Electrochem. Commun. 2007, 9, 149. (154) Kang, S.-H.; Thackeray, M. M. Electrochem. Commun. 2009, 11, 748. (155) Song, Y.; Zavalij, P. Y.; Whittingham, M. S. J. Electrochem. Soc. 2005, 152, A721. (156) Hautier, G.; Jain, A.; Ong, S. P.; Kang, B.; Moore, C.; Doe, R.; Ceder, G. Chem. Mater. 2011, 23, 3495. (157) Barker, J.; Saidi, Y.; Swoyer, J. U.S. Patent 7,901,810, March 8, 2011. (158) Ling, C.; Zhang, R.; Mizuno, F. J. Mater. Chem. A 2014, 2, 12330. (159) (a) Chen, Z.; Chena, Q.; Chena, L.; Zhang, R.; Zhou, H.; Chernova, N. A.; Whittingham, M. S. J. Electrochem. Soc. 2013, 160, A1777. (b) Chen, Z.; Chena, Q.; Wang, H.; Zhang, R.; Zhou, H.; Chen, L.; Whittingham, M. S. Electrochem. Commun. 2014, 46, 67. (160) Mba, J.-M. A.; Masquelier, C.; Suard, E.; Croguennec, L. Chem. Mater. 2012, 24, 1223. (161) Lim, S. C.; Vaughey, J. T.; Harrison, W. T. A.; Dussak, L. L.; Jacobson, A. J.; Johnson, J. W. Solid State Ionics 1996, 84, 219. 11442

dx.doi.org/10.1021/cr5003003 | Chem. Rev. 2014, 114, 11414−11443

Chemical Reviews

Review

(162) Barker, J.; Saidi, M. Y.; Swoyer, J. L. J. Electrochem. Soc. 2004, 151, A796. (163) Gaubicher, J.; Mercier, T. L.; Chaber, Y.; Angenault, J.; Quarton, M. J. Electrochem. Soc. 1999, 146, 4375. (164) Girgsdies, F.; Schneider, M.; Brueckner, A.; Ressler, T.; Schloegl, R. Solid State Sci. 2009, 11, 1258. (165) Amorós, P.; Marcos, M. D.; Roca, M.; Alamo, J.; BeltránPorter, A.; Beltrán-Porter, D. J. Phys. Chem. Solids 2001, 62, 1393. (166) Tachez, M.; Theobald, F. R.; Bordes, E. J. Solid State Chem. 1981, 40, 280. (167) Jordan, B.; Calvo, C. Can. J. Chem. 1973, 51, 2621. (168) Dupre, N.; Wallez, G.; Gaubicher, J.; Quarton, M. J. Solid State Chem. 2004, 177, 2896. (169) Girgsdies, F.; Dong, W.-S.; Bartley, J.; Hutchings, G.; Schlögl, R.; Ressler, T. Solid State Sci. 2006, 8, 807. (170) Lavrov, A. V.; Nikolaev, V. P.; Sadikov, G. G.; Porai-Koshits, M. A. Sov. Phys.-Dokl., Engl. Transl. 1982, 27, 680. (171) Bianchini, M.; Ateba-Mba, J. M.; Dagault, P.; Bogdan, E.; Carlier, D.; Suard, E.; Masquelier, C.; Croguennec, L. J. Mater. Chem. A 2014, 2, 10182. (172) (a) Azmi, B. M.; Ishihara, T.; Nishiguchi, H.; Takita, Y. J. Power Sources 2003, 119−121, 273. (b) Chen, C. Vanadyl Phosphates. Ph.D. Dissertation, Binghamton University, Binghamton, NY, 2004. (c) Allen, C. J.; Jia, Q.; Chinnasamy, C. N.; Mukerjee, S.; Abraham, K. M. J. Electrochem. Soc. 2011, 158, A1250. (173) Kerr, T. A.; Gaubischer, J.; Nazar, L. F. Electrochem. Solid-State Lett. 2000, 3, 460. (174) Harrison, K. L.; Manthiram, A. Chem. Mater. 2013, 25, 1751. (175) Barker, J.; Saidi, M. Y.; Swoyer, J. L. J. Electrochem. Soc. 2003, 150, A1394. (176) Barker, J.; Gover, R. K. B.; Burns, P.; Bryan, A.; Saidi, M. Y.; Swoyer, J. L. J. Electrochem. Soc. 2005, 152, A1776. (177) Mba, J.-M. A.; Croguennec, L.; Basir, N. I.; Barker, J.; Masquelier, C. J. Electrochem. Soc. 2012, 159, A1171. (178) Barker, J.; Gover, R. K. B.; Burns, P.; Bryan, A. Electrochem Solid-State Lett. 2005, 8, A285. (179) Ren, M. M.; Zhou, Z.; Su, L. W.; Gao, X. P. J. Power Sources 2009, 189, 786. (180) Dupré, N.; Gaubicher, J.; Mercier, T. L.; Wallez, G.; Angenault, J.; Quartron, M. Solid State Ionics 2001, 140, 209. (181) Song, J.; Xu, M.; Wang, L.; Goodenough, J. B. Chem. Commun. 2013, 49, 5280. (182) (a) Zhou, H.; Upreti, S.; Chernova, N. A.; Whittingham, M. S. Electrochem. Soc. Abstracts of the IMLB Meeting, Montreal, Canada 2010, MA2010-03, 666. (b) Zhou, W.; Upreti, S.; Chernova, N. A.; Hautier, G.; Ceder, G.; Whittingham, M. S. Chem. Mater. 2011, 23, 293. (183) Nishimura, S.-I.; Nakamura, M.; Natsui, R.; Yamada, A. J. Am. Chem. Soc. 2010, 132, 13596. (184) Tamaru, M.; Barpanda, P.; Yamada, Y.; Nishimura, S.-i.; Yamada, A. J. Mater. Chem. 2012, 12, 24526. (185) Park, C. S.; Kim, H.; Shakoor, R. A.; Yang, E.; Lim, S. Y.; Kahraman, R.; Jung, Y.; Choi, J. W. J. Am. Chem. Soc. 2013, 135, 2787. (186) Kim, H.; Lee, S.; Park, Y.-U.; Kim, H.; Kim, J.; Jeon, S.; Kang, K. Chem. Mater. 2011, 23, 3930. (187) Wen, B.; Chernova, N. A.; Zhang, R.; Wang, Q.; Omenya, F.; Fang, J.; Whittingham, M. S. Chem. Mater. 2013, 25, 3513. (188) Lister, S. E.; Rixom, V. J.; Evans, J. S. O. Chem. Mater. 2010, 22, 5279. (189) Uebou, Y.; Okada, S.; Yamaki, J.-I. J. Power Sources 2003, 115, 119. (190) (a) Ban, C.; Li, Z.; Wu, Z.; Kirkham, M. J.; Chen, L.; Jung, Y. S.; Payzant, E. A.; Yan, Y.; Whittingham, M. S.; Dillon, A. C. Adv. Energy Mater. 2011, 1, 58. (b) Li, Z.; Ban, C.; Chernova, N. A.; Wu, Z.; Upreti, S.; Dillon, A.; Whittingham, M. S. J. Power Sources 2014, 268, 106. (191) (a) Petersburg, C.; Li, Z.; Chernova, N.; Whittingham, M. S.; Alamgir, F. J. Mater. Chem. 2012, 22, 19993. (b) Sathiya, M.; Rousse, G.; Ramesha, K.; Laisa, C. P.; Vezin, H.; Sougrati, M. T.; Doublet, M.-

L.; Foix, D.; Gonbeau, D.; Walker, W.; Prakash, A. S.; Hassine, M. B.; Dupont, L.; Tarascon, J.-M. Nat. Mater. 2013, 12, 827. (c) Koga, H.; Croguennec, L.; Ménétrier, M.; Douhil, K. I.; Belin, S. E.; Bourgeois, L.; Suard, E.; Weill, F. O.; Delmas, C. J. Electrochem. Soc. 2013, 160, A786. (192) (a) Rouxel, J. Mol. Cryst. Liq. Cryst. Sci. Technol., Sect. A 1998, 310, 1. (b) Rouxel, J. Chem.Eur. J. 1996, 2, 1053. (193) Croy, J. R.; Kim, D.; Balasubramanian, M.; Gallagher, K.; Kang, S.-H.; Thackeraya, M. M. J. Electrochem. Soc. 2012, 159, A781. (194) Li, Y.; Bettge, M.; Polzin, B.; Zhu, Y.; Balasubramanian, M.; Abrahama, D. P. J. Electrochem. Soc. 2013, 160, A3006. (195) (a) Xu, B.; Fell, C. R.; Chic, M.; Meng, Y. S. Energy Environ. Sci. 2011, 4, 2223. (b) Borisevich, A. Y.; Chi, M.; Unocic, R. Electrochem. Soc. Interface 2014, 23-2, 61. (196) (a) Hong, J.; Seo, D.-H.; Kim, S.-W.; Gwon, H.; Oh, S.-T.; Kang, K. J. Mater. Chem. 2010, 20, 10179. (b) Amalraj, F.; Talianker, M.; Markovsky, B.; Sharon, D.; Luba Burlaka, A.; Shafir, G.; Zinigrad, E.; Haik, O.; Aurbach, D.; Lampert, J.; Schulz-Dobrick, M.; Garsuchc, A. J. Electrochem. Soc. 2013, 160, A324. (197) McCalla, E.; Rowe, A. W.; Camardese, J.; Dahn, J. R. Chem. Mater. 2013, 25, 2716. (198) Kam, K. C.; Mehta, A.; Heron, J. T.; Doeff, M. M. J. Electrochem. Soc. 2012, 159, A1383. (199) Noh, H.-J.; Youn, S.; Yoon, C. S.; Sun, Y.-K. J. Power Sources 2013, 233, 121. (200) Li, Z.; Chernova, N. A.; Feng, J.; Upreti, S.; Omenya, F.; Whittingham, M. S. J. Electrochem. Soc. 2012, 159, A116. (201) Kang, S.-H.; Lu, W.; Gallagher, K. G.; Park, S.-H.; Pol, V. G. J. Electrochem. Soc. 2011, 158, A936. (202) Croguennec, L.; Bains, J.; Bréger, J.; Tessier, C.; Biensan, P.; Levasseur, S.; Delmas, C. J. Electrochem. Soc. 2011, 158, A664. (203) (a) Lu, W.; Belharouak, I.; Vissers, D.; Amine, K. J. Electrochem. Soc. 2006, 153, A2147. (b) Belharouak, I.; Lu, W.; Liu, J.; Vissers, D.; Amine, K. J. Power Sources 2007, 174, 905. (c) Belharouak, I.; Lu, W.; Vissers, D.; Amine, K. Electrochem. Commun. 2006, 8, 329. (204) Zhou, F.; Zhao, X.; Bommel, A. V.; Xia, X.; Dahn, J. R. J. Electrochem. Soc. 2011, 158, A187. (205) Thackeray, M. M.; Kang, S.-H.; Johnson, C. S.; Vaughey, J. T.; Benedek, R.; Hackney, S. A. J. Mater. Chem. 2007, 17, 3112. (206) Gallagher, K. G.; Goebel, S.; Greszler, T.; Mathias, M.; Oelerich, W.; Eroglu, D.; Srinivasanae, V. Energy Environ. Sci. 2014, 7, 1555. (207) SionPower. http://www.sionpower.com/technology.html (accessed August 24, 2014).

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dx.doi.org/10.1021/cr5003003 | Chem. Rev. 2014, 114, 11414−11443