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Alexander Kraytsberg and Yair Ein-Eli *

Higher, Stronger, Better … A Review of 5 Volt Cathode Materials for Advanced Lithium-Ion Batteries

The ever-increasing demand for high-performing, economical, and safe power storage for portable electronics and electric vehicles stimulates R&D in the fi eld of chemical power sources. In the past two decades, lithium-ion technology has proven itself a most robust technology, which delivers high energy and power capabilities. At the same time, current technology requires that the energy and power capabilities of Li-ion batteries be ‘beefed up’ beyond the existing state of the art. Increasing the battery voltage is one of the ways to improve battery energy density; in Li-ion cells, the objective of current research is to develop a 5-volt cell, and at the same time to maintain high specifi c charge capacity, excellent cycling, and safety. Since current anode materials possess working potentials fairly close to the potential of a lithium metal, the focus is on the development of cathode materials. This work reviews and analyzes the current state of the art, achievements, and challenges in the fi eld of high-voltage cathode materials for Li-ion cells. Some suggestions regarding possible approaches for future development in the fi eld are also presented.

1. Introduction

The most essential parameters in chemical energy storage devices (batteries) are specifi c energy, energy density (in both cases, the larger the better), cost (the lower the better), and safety. The cell specifi c energy and energy density depend, fi rst of all, on the cell chemistry, being refl ected in its potential and charge capacity values. From this standpoint, Li-based cells hold much promise because Li metal is the most electroposi-tive (E 0 = −3.04 V vs. standard hydrogen electrode) and light ( ρ = 0.53 g cm − 3 ) material. However, employing Li metal in a secondary cell is challenging, since the possibility of dendrite growth poses risks of anode-cathode shorting (followed by the instant release of all stored energy).

In the 1970s–1980s, the concept of a Li-ion cell (“rocking chair battery”) was demonstrated; [ 1–4 ] this concept was based on the substitution of a Li metal anode with Li-ion intercala-tion compounds. The lithium is in an “almost atomic” state in a carbonaceous anode material, and it is in an “almost Li + ”-state

© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, Weinhewileyonlinelibrary.com

DOI: 10.1002/aenm.201200068

Dr. A. Kraytsberg, Prof. Y. Ein-EliDepartment of Materials EngineeringTechnion-Israel Institute of TechnologyHaifa 32000, Israel E-mail: eineli@tx.technion.ac.il

inside the cathode material, being oxi-dized by a transition metal redox couple. Whereas lithium mobility in the carbon anode is suffi ciently high, the development of cathode materials with substantial Li + -mobility turned out to be an issue of prime importance. Such a material was fi rst pre-sented by Whittingham, [ 1 ] who employed a TiS 2 -based cathode material in a cell with a metallic Li anode. The structure of TiS 2 comprises layers of hexagonal close-packed octahedral atomic groups, formed by a layer of titanium atoms between two layers of sulfur atoms, [ 5 ] thus allowing insertion of Li + into the layered gap. Upon discharge, Li + ions occupy the vacant octa-hedral sites between the layers; the charge balance is maintained by electron current via the external circuit, converting Ti 4 + into Ti 3 + . A reverse process occurs on charging, maintaining the pristine TiS 2 structure. This work promoted research on other

sulfi des and chalcogenides during the 1970s and 1980s; how-ever, cells employing such cathodes exhibited insuffi cient volt-ages of V cell < 2.5 V. In the beginning of the 1980s Goodenough et al. started working with oxide cathode material LiCoO 2 ; this layered oxide, having the structure similar to the structure of LiTiS 2 , demonstrates V cell > 4 V. [ 2 ]

The approach paved the way to safe Li-ion cells but required the development of practical anode/cathode materials, which remain undamaged over numerous Li + insertion/extraction cycles; also, the development of adequate nonaqueous electro-lytes [ 6 , 7 ] was needed. In the early 1990s, Sony succeeded in the commercialization of the fi rst rechargeable Li-ion cell based on a carbon anode (petroleum coke) and a LiCoO 2 cathode; the cell demonstrated an open circuit voltage of over 3.6 V and an energy density of ∼ 150 Wh kg − 1 . [ 8 , 9 ] Since then, Li-ion bat-teries have been recognized as high energy and high operation voltage, rechargeable power sources, [ 10 ] outperforming other available battery systems in terms of energy density, design fl exibility, cycle life, and low self-discharge rate. These features make them the ideal choice for mobile electronic devices and also an appealing option for hybrid and electric vehicle energy storage.

Current R&D in this fi eld is focused on developing high voltage cathode materials with a high charge capacity and cycling capability; substantial efforts are also being applied towards the development of organic electrolytes with a broad voltage window and high conductivity. In addition, the issues of

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Alexander Kraytsberg graduated from the Moscow Institute of Physics and Technology in 1972 and joined the Russian Academy of Science to study the electrochemistry of oxides, semiconductors, and photo-electrochemistry. During 1994–2002 he worked for Tracer Technologies Inc. (Somerville, MA, USA) con-

ducting Li-ion battery-related research. During 2003–2008 he was a Chief Scientist of Hi-Cell Ltd. (Israel) and studied DMFC technology. In 2009 he joined the Department of Materials Engineering at the Technion. Currently his work is focused on the electrochemistry of oxides, non-aqueous electrochemistry, and materials for power storage.

Yair Ein-Eli After graduating from Bar-ilan University in 1995, Prof. Ein-Eli was a post doctoral fellow at Covalent Associates Inc located at Woburn MA, USA (1995–1997), where he eventually headed the Li-ion research group until 1998. He then proceeded and joined Electric Fuel Ltd. and was appointed Director of Research and

Battery Technology. In 2001 he joined the Department of Materials Engineering at the Technion. Current research interests involved materials for batteries, solar processing and corrosion inhibitors studies.

the Li-ion material cost and their environmental compatibility now receive much attention.

2. Rechargeable Li-Ion Batteries: The Concept

2.1. Li-Ion Battery Basics

The operating principle of a Li-ion battery is illustrated in Figure 1 . The cell consists of an anode, a cathode, an electro-lyte and a separator. Lithium ions reversibly intercalate and de-intercalate into/from the anode and cathode materials on operation. The materials consist of a host material with Li + -ions accessible to inter-atomic sites. Lithium ion intercalation/de-intercalation causes a change in the charge distribution inside the host material skeleton and an overall change in the material charge which, in turn, causes electron fl ow in the external cir-cuit. Figure 1 is a schematic representation of Li-ion cell with a carbon-based anode and a metal-oxide based cathode.

The thermodynamic value of a Li-ion cell voltage (which in the absence of corrosion reactions is equal to the open circuit potential, V OC ) is determined by the difference in the chemical potentials of Li into the cathode μLi

cath and the anode μLian :

Vcell =∣∣∣∣μLi

cath − μLian

F

∣∣∣∣ (1)

Here F stands for Faraday constant.

2.2. Design Considerations

In order to fabricate a Li-ion cell with high volumetric and/or gravimetric energy density, a long cycle life, and a safe opera-tion, it has to exhibit low energy losses in course of charge/dis-charge cycle, and also a high power performance, needed for some applications. Finally, the cell must be environmentally friendly and inexpensive.

The volumetric and/or high gravimetric energy capacity is governed by the relation E stored = Q cell × V cell (here Q cell stands for the cell charge capacity); the relation suggests that enhance-ment in Li-ion cell voltage is the appropriate approach to increasing the cell energy (the charge capacity should not be

© 2012 WILEY-VCH Verlag Gm

Figure 1 . Schematics of a typical Li-ion cell.

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compromised). Since the carbonaceous anode potential (vs. Li/Li + ) Vanode

Li

/Li+ ∼ 0, the cell potential, V cell , is dictated by the

cathode potential VcathodeLi

/Li+ , i.e., by the type of cathode mate-

rial employed.

2.3. Electrolyte Requirements

Cell potential, V cell is not governed by cell electrolyte, but the employment of inadequate electrolyte compromises cell per-formance: the electrolyte should not experience oxidation/reduction at the electrode surface in the course of operation. If the electrolyte solution comprises solvent, S, and ion species S{A + }{B} − , the thermodynamic conditions of electrolyte stability are that the potentials of the electrolyte redox reactions should be in the cell voltage window, as shown in Figure 2 . The neces-sary conditions for these potentials are given by Equation 2 :

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Figure 3 . Correlation of calculated HOMO energies of organic solvents with the experimental oxidation potentials; all of the electrolytes contain 0.5 M Et 4 NBF 4 as a supporting electrolyte salt. The solvents belong to the following classes: carbonates ( � ); lactones (�); nitriles (�); formamides ( � ); oxazolidinones ( Δ ); nitroalkanes ( � ); sulfur-containing compounds ( � ). Reproduced with permission. [ 13 ] Copyright 2002, Elsevier.

Figure 2 . Electrolyte/electrode interface energy diagram in a Li-ion cell at V OC .

Electr ol yte s tabil i ty window (Wel )

=∣∣∣Vr ed

S{A+}{B−} − VoxS{A+}{B−}

∣∣∣ > Vcell

(2)

Here Vr edS{A+ }{B−} is the potential of the solvent-ion species

reduction and VoxS{A+ }{B−} is the potential of the solvent-ion spe-

cies reduction. It is convenient to infer Vred

S{A+ }{B−} and VoxS{A+ }{B−} values

by assuming that the potential for an oxidation reaction ( E ox ) correlates with the energy of the highest occupied molecular orbital ( E HOMO ) of the oxidizing species and the potential for a reduction reaction ( E red ) correlates with the energy of the lowest unoccupied molecular orbital ( E LUMO ) of the reducing species:

ELUMO = E r ed − �E sol

r + cons t1 (3)

EHOMO = E ox − �E sol

o + cons t2 (4)

Here Δ E sol r and Δ E sol o are the differences between the sol-vation energies of the pristine molecules and their ionizing forms, respectively.

Equation 2 and Equation 4 include the solvation energies, which are different for the same redox couple for different sol-vents and also different for different couples in the same sol-vent. [ 12 ] Practically, Equation 2 and Equation 4 are quasi-linear (see Figure 3 ) [ 13 ] and thus, it is possible to use theoretically cal-culated E LUMO and E HOMO as guidelines.

The computational results fi t the estimation of the redox potentials of candidate electrolytes but not for an accurate prediction of E red and E ox (and W el ) [ 13 , 14 ] and, therefore, more emphasis is to be put on experimental results for Vred

S{A+ }{B−} and Vox

S{A+ }{B−} . There is some uncertainty in an experimental determination of these values since electrochemical decomposi-tion is usually a complicated process, which is determined by both thermodynamic as well as kinetic factors. Thus, electro-chemical stability data reported for an electrolyte may depend on the conditions under which they are obtained; as a result, the literature data on W el are not always in good agreement with each other.

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The detailed consideration of the measurement approaches and of the reported values for V red

S{A+ }{B−} and VoxS{A+ }{B−} for

various electrolytes are out of the scope of this review. It is enough to mention that currently the highest reported values of W el are slightly higher than 5 V (e.g., ethylmethoxyethoxye-thyl sulfone + 1 M Li[bis(trifl uoromethanesulfonyl)imide]) [ 15 ] and, therefore, at present, the development of cathode mate-rials with Vcathode

Li

/Li+ ∼ 5 V certainly makes sense and is most

appealing.

2.4. Anode and Cathode Material Requirements

Cell potential, V cell is determining by the difference between μLi+

cath and μLi+an , and since μLi+

an > μLi

metal (preventing Li-plating), V cell may be enhanced only by increasing μLi+

cath . Since the most common carbon-based anodes weigh on average 2.5 times less than the oxide cathodes (cathode contributes ∼ 50% of the total cell weight), [ 16 ] research in the fi eld of cathode mate-rials promises more gain regarding cell energy density.

Achieving a good cell cycle life requires that electrode host compounds demonstrate a fairly good structural stability and minimal volume change over the entire operational Li-insertion/extraction voltage range.

Since Ohmic losses compromise cell effi ciency and power density, electrode materials need to maintain good electronic and Li + -ion conductivity; this problem is related mostly to cathode materials, though, because carbon has high electronic and Li + conductivity.

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Figure 4 . Schematic representation of a band structure of a Li-insertion material; slightly oxidized redox couple as it passes through the top of the p an : (a) itinerant vs. polaronic character of holes states of couple on the approach to the top of p an , (b) pinned couple with predominantly p an holes states. Adapted with permission. [ 19 ] Copyright 2009, Elsevier.

Figure 5 . Schematic representation of a layered LiMO 2 structure (M stands for Co, Ni, or Mn). Adapted with permission. [ 24 ] Copyright 2007, Elsevier.

3. Lithium Transition Metal Oxide-Based Cathodes

3.1. Decomposition Limit of the Cathode Material

The potential-determining reaction of a metal-oxide cathode may be represented as δ Li + + Li x M y O z DLi x + δ M y O z , and the potential μcath

x depends on the degree of lithiation, x . At the same time, there is a possibility that an oxide is involved in a parallel reaction of decomposition: V cath

x0 .

At a certain value of V cathx0

, i.e., at lithiation x 0 , starting from V cath

x0 the cathode experiences decomposition instead

of charging; the decomposition onset potential V cathx0

is an intrinsic cathode material parameter. It is reported that, in the course of de-lithiation of Li 1 − x CoO 2 where x > 0.5, electrons are transferred from the oxygen band p an , and O 2 evolves. [ 17 ]

Oxidative decomposition of the metal oxide may be discussed based on the concept of pinning of a transition metal redox energy level. [ 18 , 19 ] Li-insertion compounds may be considered semiconductors with a conductive band ( CB ) formed by tran-sition metal d-orbitals ( d met ) and a valence band ( VB ) formed by oxygen (or sulfur, etc.) p-orbitals ( p an ); in a real material the metal–anion bonding is not 100% ionic, thus CB is formed by hybrid orbitals ( d met ∪p an ), with a dominant input of d met , and VB is formed with a dominant input of p an (see Figure 4 ).

As the material oxidation increases (i.e., as the charge carrier concentration changes), the redox level along with the Fermi level E F moves toward VB and the p an share in the hybridiza-tion increases. Finally, the transition metal redox couple-related band reaches the ceiling of VB (Figure 4 b), and E F pinning takes place. Starting from this point, if further Li + de-intercalation follows, electrons tend to be taken from the ceiling of VB ; i.e., holes appear in VB , forming di-anionic states like (O 2 ) 2 − , followed by a disproportionation of these states according to

© 2012 WILEY-VCH Verlag GmbAdv. Energy Mater. 2012, 2, 922–939

Equations 1 and 2 and a fi nal degradation of the cathode mate-rial following Equation 5 is expected (more elaborated discus-sions of this approach may be found in the literature [ 19 , 20 ] ):

2 (O2)2− → 2O2− + O2 ↑ (5)

This approach suggests that V cathx0

may depend on fi ne details of the oxide cathode electronic structure. Indeed, a removal of Li + from Li x M y O z can be considered as the addition of a hole. The hole may be located mostly in oxygen (which favors oxide decomposition) or it may be distributed over both the metal and oxygen. For example, Li x CoO 2 starts to decompose at x < 0.5 [ 17 ] whereas Li x NiO 2 may be completely de-lithiated (down to NiO 2 ) without any decomposition. [ 21 ] At the present, ab-initio quantum mechanics calculations are not able to provide an accurate expla-nation of the hole distribution upon de-lithiation [ 22 ] and, therefore, only experimental data on V cath

x0 may be considered reliable.

3.2. Major Types of Compounds Employed

As outlined above, electrode materials for Li-ion cell cathode should maintain a good Li + -ion conductivity; currently, three types of metal oxides are used as cathode materials in lithium-ion batteries; [ 23 ]

• Layered oxides with the α -NaFeO 2 -type structure • Oxides with a spinel structure, with (Fd 3 m) symmetry • Poly-anion oxides with the olivine and olivine-related

structures, with (Pnma) symmetry

These types of the materials are considered and discussed in the following sections, focusing on the development of high voltage cathodes.

3.3. Layered Oxides

These oxides with the general formula of LiMO 2 form a distorted rock-salt ( α -NaFeO 2 -type) crystal structure. All known structures are derived from anionic stacking: 03, P3, 02, and P2. Here the let-ters O and P stand for the Li site (octahedral or prismatic) and the number indicates the amount of layers contained in the elementary crystallographic cell; the < MO 2 > -layers are built up of edge-sharing < MO 6 > octahedra ( Figure 5 ). These layers are not deconstructed

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Figure 7 . Cycle characteristics of LiNiO 2 under a constant-charge capacity, 1.0 mA cm 2 , 1 M LiCIO 4 in a solution of propylene carbonate and dimeth-oxyethane. Adapted with permission. [ 28 ] Copyright 1995, Elsevier.

during Li + intercalation/de-intercalation. The oxygen ions form a cubic-close-packed array along the C axis and the Li + and M 3 + ions alternately occupy the close-packed (111) interstitial octahedral-site planes. Two-dimensional Li + ion diffusion in the space between < MO 2 > layers occurs since the edge-shared octahedral < LiO 6 > arrangement allows ion movement from one (vacant) octahedral-site plane to another via a tetrahedral site. The edge sharing of < MO 6 > octahedra maintains a direct < M–M > interaction, respon-sible for the material’s electronic conductivity. [ 24 ] Thus, layered metal oxides are the most commonly used cathode materials. [ 25 ]

3.3.1. LiCoO 2

Li x CoO 2 [ 25 ] is the most popular cathode material with fair elec-trical conductivity and Li + mobility. Consistent with the high redox value of the Co 4 + /Co 3 + couple, its operational voltage is around 4 V. [ 2 , 3 , 25 ] A capacity of 274 mAh/g corresponds to a full Li + extraction, producing CoO 2 , which generally is possible. [ 26 ] However, a reversible extraction/insertion of Li + is practically limited to x = ∼ 0.5, yielding a capacity of 140 mAh/g; the mate-rial demonstrates hundreds of cycles within this range, but de-intercalation of lithium below x = 0.5 results in a substantial capacity fade, which is attributed to material decomposition. It manifests as oxygen evolution, [ 17 ] cobalt dissolution, and a serious lattice shrinkage: on full Li loss, inter- < CoO 2 > -layer dis-tance shrinks from ∼ 1.422 down to 1.288 nm, [ 26 ] which is large enough to cause material crumbling. Major material drawbacks are the high cost of cobalt and its environmental hazard.

3.3.2. LiNiO 2

Nickel is fairly inexpensive and is also a fairly environmen-tally friendly element. Thus, it is not a surprise that LiNiO 2 (its structure is a slightly Jahn-Teller distorted version of α -NaFeO 2 ) had attracted much of an attention. It has a similar operation potential (somewhat lower than LiCoO 2 , though, namely, 4.1 V for Li ½ CoO 2 vs. 3.95 V for Li ½ NiO 2 ) [ 27 , 28 ] but develops higher sta-bility on de-lithiation than LiCoO 2 ; it may be even de-lithiated to form NiO 2 [ 29 ] and thus, it may be cycled up to 200 mAh/g. [ 28 ] The most interesting feature is that the NiO 2 Li + -intercalation voltage is 4.8 V, and electrons are extracted from the e g band. Since the e g band lies well above the O2p band, lattice oxygen is not displaced even at a high degree of de-lithiation ( Figure 6 ). [ 21 ]

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Figure 6 . Energy diagrams of Li 1−x CoO 2 and Li 1−x NiO 2 . Adapted with per-mission. [ 21 ] Copyright 2001, American Chemical Society.

However, LiNiO 2 employment meets some problems (and hence, challenges) for a practical Li-ion cell. The fi rst problem is associated with the preparation of well-ordered LiNiO 2 . In course of the preparation, Ni 2 + /Li + -containing precursors must form a new compound so that Ni 2 + is oxidized and both ele-ments have to settle in their new places to form a LiNiO 2 com-pound with an α -NaFeO 2 structure. Nickel ions are only suffi -ciently itinerant in the Ni 2 + /Li + -oxide precursor mix at tempera-tures above 600 ° C, [ 30 ] whereas Ni + 3 -ions are not stable at high temperatures. As a result, up to now experimentalists have failed to avoid the presence of divalent nickel ions, and half of them occupy the lithium sites between < NiO 2 > layer space. [ 31 ] Being larger than Li + (0.83 Å vs. 0.74 Å), [ 32 ] Ni + 2 substantially reduces Li + mobility and thus degrades the electrode power capability of the material; in fact, the Li + -diffusion coeffi cient in commonly prepared LiNiO 2 is several times less than that in LiCoO 2 . [ 33 ]

The second problem is that LiNiO 2 actually has several crystal modifi cations, the most electrochemically favorable being LiNiO 2 (R 3 m, α -NaFeO 2 -type phase), with a layered structure, and (among others) LiNiO 2 (Fm3m), with a rock salt-type structure; the lattice parameters of both modifi cations are very close and thus the modifi cations coexist, but the latter structure is electrochemically inactive. As a result, commonly prepared LiNiO 2 (R 3 m) is contaminated with electrochemically inactive rock-salt domains. Numerous synthetic routes have been offered to bypass these obstacles, [ 31 , 34 ] though an adequate solution has not yet been found.

The third problem is that, upon de-lithiation, Li 1− x NiO 2 exhibits several successive phase transformations accompa-nied by substantial volume changes; this seriously degrades the material integrity and compromises the cycling ability. The result is that Li 1− x NiO 2 may be cycled only in a consid-erably narrow interval of x ( Figure 7 ) [ 28 ] and, therefore, the high-voltage capability of the material cannot be exploited. An additional problem is the low thermal stability of the low-Li

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Figure 8 . Thermogravimetric analysis-based data on Li 0.3 NiO 2 , Li 0.4 CoO 2 , and λ -MnO 2 ; oxygen release onset temperature t onset is presented as a function of the heating rate V heat , i.e., as t onset (V heat ); t onset is determined by t onset (V heat ) extrapolation to V heat = 0. Adapted with permission. [ 36 ] Copyright 1994, Elsevier.

Li x NiO 2 ( x � � 1), which contains unstable Ni 4 + ions. Ni + 4 may be reduced back to Ni 3 + at elevated temperatures; the process Ni 3 + → Ni + 4 results in a change to the lattice structure (spe-cifi cally, α -NaFeO 2 -type phase transfers into a pseudo-spinel phase and then into a highly disordered R3hm phase), which is accompanied by oxygen loss, as shown in Equation 6; the smaller the value of x , the lower the Li x NiO 2 -phase temperature stability: [ 35 ]

Li0.30Ni1.02O2 → 0.12O2 + Li0.30Ni1.02O1.76 (6)

This possibility of oxygen release seriously compromises cell safety; its onset temperature value t onset is crucial: the loss of oxygen will not compromise battery safety if the t onset is higher than the practical operation temperature. It has been demon-strated ( Figure 8 ) that this Li 0.3 NiO 2 t onset is just above 100 ° C (though the decomposition is fairly slow up to 150 ° C). [ 36 ]

Numerous attempts to improve the properties of Ni-based materials have been, up to now, mostly focused on substituting some other ion for nickel; the fi rst candidate is (as expected) cobalt. The LiNi 1−y Co y O 2 compounds and adequate synthetic routes to them have been extensively studied; the presence of Co instead of Ni can stabilize the layered structure and sup-press cation disorder, enabling a reversible capacity close to 180 mAh/g; for the optimal formulation, y should be within the interval 0.15–0.3. Cobalt also suppresses the phase transitions associated with Li + extraction from LiNiO 2 . Whereas low cost and fair capacity makes LiNi 1−y Co y O 2 compounds attractive can-didates for a cathode material, some problems are yet unsolved, namely, capacity fade during cycling, which is due to the migra-tion of Ni 3 + ions from the Ni planes to the Li planes still being too high (LiCoO 2 has better cyclability). The operating voltages lie in the range between those of LiNiO 2 and LiCoO 2 materials, and safety issues, which are related to the possibility of oxygen

© 2012 WILEY-VCH Verlag GmAdv. Energy Mater. 2012, 2, 922–939

gas formation during cell operation, persist. [ 24 , 37 , 38 ] Additives such as Al, Mg, Ca and Ba have also been employed. The intro-duction of these ions results in the elimination of phase trans-formations and, thus, diminishes material volume changes during cycling. The thermal stability of LiNiO 2 and LiNi 1−y Co y O 2 may also be enhanced by Al, Mg, and Mn doping. [ 24 , 39–45 ] The bottom line is, though, that all these LiNiO 2 -based doped mate-rials generally show lower operating voltages and lower cycling capabilities than those of LiCoO 2 .

3.3.3. LiMnO 2

Thia material looks quite attractive, since Mn is inexpensive and more environmentally friendly than Co or Ni. The stable phase of LiMnO 2 has an orthorhombic structure, [ 46 ] and layered LiMnO 2 is metastable; layered LiMnO 2 provides easy Li + -diffusion paths, while also exhibiting a fairly smooth intercalation/de-intercalation voltage profi le. [ 11 ] However, although most Li ions can be removed from the structure on charging (providing a capacity of ∼ 285 mAh/g), the following Li + insertion does not restore the layered LiMnO 2 . On extraction of half the Li from the layered LiMnO 2 , manganese ions penetrate the interlayer space, and this results in the formation of crystal regions with a spinel structure. [ 11 , 47 , 48 ] This phase transformation is irreversible and compromises the cathode cycle life. The operation voltages of the layered LiMnO 2 areall found to be lower than 4.1 volts, [ 11 ] which is consistent with quantum-mechanical calculations. [ 49 ]

Attempts to stabilize layered LiMnO 2 by doping the com-pound with elements such as Co, Ni, Cr, etc., are the subject of several experimental and theoretical studies. [ 48 , 50–52 ] It was reported that Li(Ni 1/3 Mn 1/3 Co 1/3 )O 2 can operate at voltage of 4.5 V with a capacity of around 200 mAh/g; the cyclability was not reported. [ 53 ] Nevertheless, these studies have only a limited success and the cyclability is still an issue. More important, from the point of the subject of the present review, the working voltages of modifi ed LiMnO 2 are still below 4.5 V, and the bulk of their capacity is associated with voltages even below 4 V.

3.3.4. On the Chance to Develop a Layered Oxide Material with an Elevated Voltage

Whereas current ab-initio quantum-mechanical computations cannot deliver the exact values of V cath

x , these computations are useful for material pre-screening before their synthesis and investigation. The computations also are quite helpful in revealing the detailed mechanism of the lithium intercalation process. [ 22 , 54 , 55 ] Particularly, it was revealed that the intercalation voltage is governed not only by the redox potential of the transi-tion-metal ion, which changes its valency upon Li + intercalation; the lattice oxygen also may be involved in the electron exchange ( Table 1 ). [ 55 ] It was suggested that this increased oxygen partici-pation correlates with increased V cath

x . [ 54 ] This approach implies that it may be possible to preserve the electrochemical activity and deliver higher V cath

x substituting a part (or all) of the oxide transition metal with non-transition-metal ions.

Since oxygen participation in charge transfer is highest in the case of aluminum, it is expected that the calculated V cath

0 for LiAlO 2 will be as high as 5.4 V. [ 54 ] Pure LiAlO 2 is not electronic conductor but the mixed oxides Li(Al x Co y )O 2 have the α -NaFeO 2

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Copyright 1997, American Physical Society.

Bond Ti-O V-O Mn-O Co-O Ni-O Cu-O Zn-O Al-O

Oxygen 0.216 0.245 0.254 0.25 0.255 0.261 0.273 0.322

Metal 0.262 0.253 0.289 0.276 0.126 0.102 0.067 0.013

Figure 10 . LiMn 2 O 4 spinel structure representation; LiMn 2 O 4 shown being composed of dark gray MnO 6 octahedra and lithium ions (light gray balls) occupying interconnected tetrahedral positions. Adapted with permission. [ 24 ] Copyright 2007, Elsevier.

structure and are semiconductors, and calculated values V cath0

for such oxides are higher than for LiCoO 2 . The experimental results are ambiguous on Al-substituted layered oxides, though, and seemingly depend on the material morphology and prepa-ration methods. Specifi cally, whereas the results of Jang et al. [ 56 ] ( Figure 9 ) support these theoretical predictions, other works do not confi rm that Al-doping causes Li + -intercalation voltage enhancement. [ 57 , 58 ] This approach paves the way to the develop-ment of new high-voltage layered oxides.

3.4. Spinel-Type Oxide Materials

Spinel oxides with the general formula LiM 2 O 4 comprise octahedral-coordinated M-cations and Li-cations in tetrahedral positions of a cubic-closed-packed O 2 − lattice. The M cations are forming a 3D-[M 2 ]O 4 framework in which the interstitial space is formed by edge-sharing octahedral sites that share faces with the tetrahedral Li-containing sites, as shown in Figure 10 . The frame-work remains stable while Li + ions reversibly move between the tetrahedral sites, and good Li + conductivity is imparted by an edge-shared < MO 6 > -octahedral arrangement with a direct < M–M > interaction: interconnected tetrahedral sites provide 3D paths for Li + diffusion through the spinel framework, which makes spinel-type oxides promising materials for Li-ion cell cathodes. [ 4 ]

Ab-initio quantum-mechanical calculations hint that spinel materials may have a higher voltage at low Li + content compared with layered materials. The calculated voltage of a virtual spinel modifi cation of LiCoO 2 is substantially higher that the voltage of layered LiCoO 2 [ 59 ] ( Figure 11 ); it was suggested that spinel

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Figure 9 . Voltage curves of Li 1−x Al y Co 1−y O 2 in charging up to x = 0.4, fol-lowed by discharging to x = 0.2, 0.4 mA cm − 2 ; (a) y = 0, (b) y = 0.25. Adapted with permission. [ 56 ] Copyright 1999, The Electrochemical Society.

Figure 11 . Comparison of calculated voltage intercalation curves for vir-tual spinel Li x CoO 2 (dashed line) and actual layered Li x CoO 2 (solid line) at 300 K. Adapted with permission. [ 59 ] Copyright 1999, American Physical Society.

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Figure 12 . Open-circuit voltage curve of spinel-related Li x Mn 2 0 4 at 30 ° C. Adapted with permission. [ 64 ] Copyright 1990, The Electrochemical Society.

Figure 13 . Variation in the Mn-O COOP as a function of the substituting element; the substitution corresponds to 25% of the entire manganese atoms. Adapted with permission. [ 68 ] Copyright 2000, The Electrochemical Society.

LiCoO 2 can be obtained experimentally but its intercalation potential was found to be substantially lower than the calcu-lated value. [ 60 ]

3.4.1. LiMn 2 O 4

It is expected that spinel LiMn 2 O 4 attracts a large amount of attention, as the combination of the spinel 3D lattice with chem-ical stability of the Mn + 3 /Mn + 4 couple promises good safety and high power capability for a cathode material. Manganese is also a fairly cost-effective material. [ 61 ] It surfaced, though, that LiMn 2 O 4 is a problematic cathode material because of its high capacity fade, especially at temperatures over 50 ° C [ 62 ] (the material loses Mn ions in the course of cycling), which makes its implementation impractical. The material may also undergo a phase transition in the course of Li de-intercalation, which compromises cycle life [ 63 ] and results in the existence of a voltage step, and even the maximal de-lithiation voltage is just a bit higher than 4 V ( Figure 12 ). [ 64 ] Therefore, a pure LiMn 2 O 4 may hardly be considered a high-voltage cathode material. Spinel-type oxides LiM 2 O 4 with M = Ti, V, Co also demonstrate the operational voltages below 4.0 V. [ 65 ]

3.4.2. LiMn 2 O 4 -Based High-Voltage Spinel Compounds

In pure LiMn 2 O 4 , manganese ions can be partially replaced by other metal ions; computations have revealed that the introduc-tion of alien cations doesn’t compromise the spinel structure [ 66 ] but modifi es the electronic properties of the material. Particu-larly, such substitution changes the crystal orbital overlap pop-ulation (COOP), [ 67 ] as demonstrated in Figure 13 , [ 68 ] and also delays the onset of Jahn–Teller distortions of the [Mn 3 + O 6 ] octa-hedra during discharge; [ 69–71 ] the latter process provokes a struc-tural transition accompanied by a volume increase, [ 72 ] which triggers cathode material crumbling. Since doping suppresses

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such structural transitions, [ 73 ] mixed ternary Li( A M) α Mn 2− α O 4 (and quaternary Li( A M) α ( B M) β Mn 2- α - β O 4 ) spinels (here B M and A M are di- and/or trivalent metal cations) may have better cyclability and, particularly, higher operating voltages than LiMn 2 O 4 . [ 74–76 ] Indeed, the Mn ion has a formal oxidation state of [Mn 3.5 + ] in spinel, and the substitution of Mn with X M metals, whose oxidation state is less than [ X M + 3.5 ], increases the average oxidation state of the remaining manganese. It is also known that the higher the average oxidation state of Mn in the material, the higher the material’s cyclability; [ 77–79 ] this is par-ticularly true for near-surface Mn. [ 80 ]

Increasing the average Mn ion valency (which accompanies doping) not only suppresses Jahn–Teller distortions and phase transitions, but also may decrease Mn dissolution. [ 80 , 81 ] It was found that Mn + 2 dissolution does not always account for the major part of the cathode material degradation. Specifi cally, it was estimated that Mn dissolution may account for less than 30% of capacity loss. [ 77 ] The reduction of Mn dissolution is important for preventing the whole Li-ion cell degradation. The associated problem with the anode is that the dissolved Mn cat-ions are deposited on the graphite anode, and this deposited Mn causes anode capacity loss. [ 82 , 83 ] The problem is that the cathode material particles dissolve, resulting in loss of contact between active material particles and conductive additives and, thus, the demise of cathode conductivity. [ 84 ] Concluding the overview of the Mn-spinel cathode material dissolution issue, it also worth outlining that dissolution depends not only on the material fea-tures but also strongly on the electrolyte content. [ 85 , 86 ]

At the same time, lower-valence substitution decreases the number of Li + ions that can be extracted from the spinel

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Figure 14 . Calculated Li + -diffusion barriers in doped LiMn 2 O 4 (LiM 0.5 Mn 1.5 O 4 ); here M stands for Cr, Fe, Co, Ni, Cu, Mn; �- one Mn-ion is substituted in 16c site–surrounding six Mn- ring, � –three Mn-ions are substituted in 16c site - surrounding six Mn-ions ring, � - non-doped LiMn 2 O 4 . Adapted with permission. [ 90 ] Copyright 2011, American Chem-ical Society. Datum on LiCoO 2 acquired and replotted. [ 89 ]

structure before all Mn ions are in the oxidation state [Mn + 4 ], which results in capacity reduction. Thus, increasing opera-tional voltage and cyclability via manganese substitution is inevitably a capacity trade-off.

The work in this fi eld is also supported by calculations of Li -mobility (cathode materials with low Li + conductivity are not practical); the calculations hint that Mn substitution may preserve Li + mobility in the doped spinels. Indeed, the spi-nels belong to the d3mF space group with oxygen ions in 32e sites forming a close-packed fcc lattice. In manganese spinel, LiMn 2 O 4 , Mn ions are located in the 16d octahedral sites, while Li ions occupy the 8a tetrahedral sites, and the 16c octahe-dral sites are empty. Lithium ion diffusion occurs via hopping between adjacent 8a sites through the intermediate 16c sites; these 16c sites are actually surrounded by six Mn ions forming a kind of Mn fence, or Mn ring. [ 87 ] Thus, Li + mobility and hence ionic conductivity (which determines material power perform-ance and energy losses) are functions of the barrier between two adjacent 8a sites. In connection with this, the higher Li + mobility and ionic conductivity of the layered oxides (if com-pared with spinel-type oxides) [ 88 ] stems from the fact that in lay-ered oxides this barrier is substantially lower. [ 89 ] Figure 14 dem-onstrates that this hopping barrier often doesn’t substantially alter Mn-spinel doping; moreover, some dopants may even increase the ionic conductivity of the material. [ 90 ]

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Experimentally, a substantial number of substituted Mn-spinels Li( A M) α ( B M) β Mn 2- α - β O 4 , which operate over 4.5 V, have been identifi ed. Among these compounds are those substituted with Ti (LiNi 0.5 Ti α Mn 1.5- α O 4 , [ 91–93 ] (LiTi α Mn 2- α O 4 doesn’t demonstrate a high-voltage plateau)), [ 94 , 95 ] with Cr (LiCr α Mn 2- α O 4 [ 74–76 , 94–96 ] and LiMn 1.5 − α Cr 2 α Ni 0.5 − α O 4 ), [ 97–100 ] with Fe, [ 75 , 76 , 101 , 102 ] with Co, [ 74 – 76 , 103 , 104 ] with Ni, [ 74 – 76 , 105–111 ] with Co and Ni, [ 98 , 112 , 113 ] with Fe and Ni, [ 106 , 114 ] with Co, Ni, and Fe (LiNi z Co x Fe y Mn 2-x-y-z O 4 ), [ 115 ] with Cu, [ 75 , 76 , 116–118 ] with Cu and Ni, [ 74 , 106 ] with Zn, [ 74 , 119 ] with Zr, [ 100 ] with Mg, [ 120 ] with Mg and Ni, [ 106 , 121 ] and with Al. [ 74 , 106 , 100 ]

3.5. High Voltage Capacity: Spinel-type Oxides vs. Layered Oxides

Up to now, most of the research in the fi eld of spinel-type mate-rials has been focused on mixed manganese-oxide based mate-rials with a general formula of Li x M y Mn 2−y O 4 (M = Ni, Co, Fe, Cr, etc.). Among these materials, there are a substantial number of cathodes with the highest de-lithiation potential, over 4.8 V and even over 5 V (e.g., see Table 2 ), [ 66 ] whereas the popular lay-ered oxides have substantially lower maximal operational volt-ages of 3.95 V for Li ½ NiO 2 [ 28 ] and 4.1 V for Li ½ CoO 2 . [ 27 ]

The issue is, however, that the discharge curves of most of these materials comprise two steps: in plain terms, the low voltage step is related to the oxidation of Mn ions, and the high voltage step is related to the oxidation of the other metal in the Li x M y Mn 2−y O 4 spinels. Thus, the high-voltage step occu-pies only part of the lithiation curve and the average discharge voltage is not too high, despite the impressive 5 V height of the voltage plateau. The 5 V-step increases the overall energy of Li x Co 0.4 Mn 1.6 O 4 only by ∼ 6%: this behavior is demonstrated by most spinels of Li-Mn-M-O systems with M = Co, Cr, Co, Fe ( Figure 15 ; and references in Table 2 ). [ 75 , 128 ] The only known exception (and one of the most studied spinel materials) is LiNi 0.5 Mn 1.5 O 4 , which, being properly prepared, may offer a one-step discharge curve. [ 105 , 111 , 129 , 130 ]

Another challenge is that most of the spinel-type oxides develop two-phase behavior; judging from the results of ab-initio computations, [ 59 ] existence of this two-phase region is not determined by a specifi c component (e.g., manganese) but is a feature of the spinel structure. The computations indicate that not only real spinel cathode materials, but also a simulated non-existent “spinel-type” single phase Li x CoO 2 , break down into two phases upon lithiation, forming a two-phase region which results in a step-wise charge curve ( Figure 16 ).

This is an unfavorable situation since the coexistence of two phases in a substantially wide interval of lithium content results in phase bordering and phase interface movement through cathode material grains; this is exactly the process which forms intergrain stresses and is detrimental for the preservation of grain integrity. At the same time, layered oxides may keep the same structure over the course of lithiation and thus avoid phase border shifts through material grains, and therefore it is not a surprise that up to now layered oxides develop better cycle life then spinels (e.g., see data on spinel-type cathodes). [ 106 ]

While considering safety issues, the most employed layered material (LiCoO 2 ) has developed decomposition with oxygen

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Table 2 . Electrochemical data for high-voltage lithium cathode materials, Li 2 MMn 3 O 8 and LiMMnO 4 based on spinel-structure oxides. Adapted with permission. [ 66 ] Copyright 2007, Elsevier.

Material Mid-discharge voltage [V] plateau > 4.5 V

Redox couple b) at plateau > 4.5 V

Plateau centered at 4.0 V [mAh g − 1 ]

Plateau > 4.5 V [mAh g − 1 ]

V range [V]

Reference

Li 2 CrMn 3 O 8 4.8 Cr 3 + /4 + 70 55 3.4–5.4 [ 122 , 123 ]

LiCrMnO 4 4.8 Cr 3 + /4 + 0 c) 75 3.4–5.4 [ 122 , 123 ]

Li 2 FeMn 3 O 8 4.9 Fe 3 + /4 + 75 50 3.0–5.3 [ 124 ]

Li 2 CoMn 3 O 8 5.1 Co 3 + /4 + 70 60 3.0–5.3 [ 104 , 66 ]

LiCoMnO 4 5.0 Co 3 + /4 + 10 95 3.0–5.3 [ 125 ]

Li 2 NiMn 3 O 8 4.7 Ni 2 + /4 + 16 95 3.0–4.9 [ 105 , 126 , 127 ]

Li 2.02 Cu 0.64 Mn 3.34 O 8 a) 4.9 Cu 2 + /3 + 48 23 3.3–5.1 [ 117 ]

a) Composition of the spinel component at the nominal composition Li 2 CuMn 3 O 8 ; b) all redox couples are located in octahedral sites; c) an inclined single plateau ranging 3.8–4.8 V was observed.

Figure 15 . Comparison of charge (100 mcA/cm 2 ) curves for Li/Li + /Li y M x Mn 2-x O 4 (M = Cr, Fe,Co, Ni and Cu) cells. Adapted with permission. [ 75 ] Copyright 2004, Elsevier.

evolution starting from T dec = 230 ° C [ 131 ] (naturally, reaction with the electrolyte may start at 190 ° C [ 132 ] or even as low as at 155 ° C). [ 133 , 134 ] The introduction of doping and alloying elements may drastically shift up the decomposition onset temperature T dec [ 135 , 136 ] (e.g., T dec ∼ 300 ° C for LiNi 0.225 Co 0.55 Mn 0.225 O 2 ) [ 136 ] and also avoid high-temperature reactions of the material with the electrolyte. [ 137 ] Aluminum (a high-voltage dopant for layered oxides) may enhance the thermal stability of layered cathode compounds or, at least, will not compromise it. [ 35 , 58 , 138–141 ] There is also the option to coat the grains of the base cathode mate-rial with a thin layer of thermally stable Li + -conducting com-pounds; this approach may increase T dec by up to 50 ° C com-pared with the uncoated material. [ 142–144 ] One more option is to

© 2012 WILEY-VCH Verlag GmbH & Co. KGaA, WeinAdv. Energy Mater. 2012, 2, 922–939

introduce fl uorine; e.g., T dec is ∼ 20 ° C higher for Li[Ni 1/3 Co 1/3 Mn (1/3-0.04) Mg 0.04 ]O 1.92 F 0.04 than for Li[Ni 1/3 Co 1/3 Mn (1/3-0.04) Mg 0.04 ]O 2 . [ 145 ] Spinel-type materials are thermally more stable; e.g., LiMn 2 O 4 decomposes with oxygen evolution starting at about 600 ° C; [ 146 ] Li x Mn 2 O 4 with x � 1 decomposes at lower temperatures, though, and the reaction with the electrolyte also takes place at lower temperatures. [ 147 ] Nickel doping decreases the decomposition temperature, shifting it as low as 240 ° C. [ 148 ] This is particularly true for the spinel high-voltage champion LiNi 0.5 Mn 1.5 O 4 [ 149 ] but its thermal stability may be substantially increased by fl uorina-tion [ 148 , 150 ] and also Cr doping. [ 151 ]

Summing up, it might be concluded that spinel compounds look more promising as candidates for commercial high voltage cathode materials (compared to layered oxides) from the points of view of intercala-tion potentials, rate capability, energy density, safety, and environmental compatibility.

Considering further improvements of spinel-based materials, the major fi eld of activity is increasing its cycle life. [ 152 ] Recently it was demonstrated that Ru-doped

LiNi 0.5 Mn 1.5 O 4 materials have an excellent cycle life and power capability, [ 153 ] and don’t develop different immiscible phases. It is now commonly accepted that a material’s particle size plays an important role in electrochemical performance [ 154 , 155 ] and specifi -cally in determining the charge/discharge voltage of the cathode materials. [ 156 ] There have been some encouraging results on grain size-related properties of spinel oxide cathodes ( Figure 17 ). [ 153 , 157 ]

3.6. Poly-Anion Compounds

Poly-anion cathode materials have attracted quite a large amount of attention following the pioneering works of

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Figure 16 . Calculated phase diagram of simulated spinel Li x CoO 2 (Co ions remain in the 16d sites). The full line bounds the region with two immiscible phases (between x = 0.5 and x = 1.0), the dashed line bounds the region where a second-order transition (disordered phase) → (ordered phase) takes place (centered on x = ¼); in the ordered phase, Li + occupy every second [8a] site while other tetrahedral and octahedral sites are empty. Adapted with permission. [ 59 ] Copyright 1999, American Physical Society.

Goodenough et al. [ 158 , 159 ] These materials include compounds with a NASICON-type crystal lattice Li x

.M2 (

.X O 4 ) 3 (here

.M is Ni,

Co, Mn, Fe, Ti or V and X. is S, P, As, Mo or W), an olivine-type

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Figure 17 . Discharge curves of LiMn 1.5 Ni 0.5 O 4 cathodes at various C-rateChemical Society.

3

3.2

3.4

3.6

3.8

4

4.2

4.4

4.6

4.8

5

0 20 40 60

Vol

ts v

s. L

i/Li

+

mAh g-1

LiNi0

crystal lattice Li MX x O 4 (here M is Fe, Co, Mn or Ni and X is P, Mo, W or S), [ 160 ] and poly-anion materials with tavorite-related structures (tavorite, triplite, maxwellite, sillimanite), with the general formula LiM M 1 − δ δ (ZO 4 )X 1− α

.X α , where at least

one of .

M or M is a metal with several possible oxidation states, Z is commonly phosphorus or sulfur, and X and

.X are oxygen,

a hydroxyl group, or a halogen (commonly fl uorine). [ 161–164 ] All these are materials with open 3D frameworks, which

are available for Li migration. These poly-anion compounds have been extensively investigated in the last decade; their cyclability, safety (the materials are less prone to decomposi-tion and releasing oxygen at elevated temperatures than corre-sponding—i.e., with the same transition metal—layered oxides and spinels), [ 134 , 165 ] environmental compatibility, and potentially low production cost make them the most promising electrode material in prospective Li-ion batteries.

NASICON frameworks are formed by .

M 2 ( .

X O 4 ) 3 ; in these frameworks all [

.M O 6 ]-octahedral corners are common with

.X

O 4 -tetrahedral corners, and all .

X O 4 -tetrahedral corners are common with [

.M O 6 ]-octahedral corners; materials with such

frameworks are known to be available for topotactic inser-tion/extraction of alkaline ions, [ 166 ] and specifi cally Li + ions, due to the availability of the open 3D framework for easy Li + movement along conducting channels. [ 167 ] This feature makes NASICONs promising materials for rechargeable lithium bat-teries. [ 168 ] The conduction pathways, which NASICON frame-works offer for Li + motion, are commonly considered 1D paths along the C -axis. [ 160 , 169 ] The conducting channels are not really straight, linear pathways, but zigzag, wavy pathways. [ 170 , 171 ] The

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s and grain sizes. Adapted with permission. [ 153 ] Copyright 2011, American

80 100 120 140

.5Mn1.5O4

1C, 300nm

1C, 1-0.6 mcm

5C, 300nm

5C, 1-0.6 mcm

10C, 300 nm

10C, 1-0.6 mcm

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Figure 18 . LiFePO 4 (olivine structure) representation; the oxide shown being comprised of dark gray FeO 6 octahedra and light gray PO 4 tetra-hedra lithium ions are shown as light gray balls occupying octahedral sites, which form linear chains of edge-shared octahedra; Reproduced with permission. [ 24 ] Copyright 2007, Elsevier.

pathway topology depends on the morphology of the material. Specifi cally, Li + hopping between pathways is possible in the presence of

.X O 4 -tetrahedral disorder, [ 172 ] which actually turns

conductivity lines into conductivity branches; it may be con-cluded that whereas Li + motion is strongly anisotropic it is not purely a 1D process and has a quasi-2D character. NASICON .

M 2 ( .

X O 4 ) 3 frameworks are versatile from the chemical point of view: their structure could accommodate a variety of transition metal cations with a set of redox potentials, and various

.X O 4

anions. The Li x MX O 4 [ M = Fe, Co, Mn or Ni and X = P, Mo, W or

S] olivine structure has Li and M atoms in the octahedral sites and X atoms in the tetrahedral sites of a hexagonal-close-packed oxygen array. With Li in a continuous chain of edge-shared octahedra on alternate planes, a reversible topotactic insertion/extraction of lithium from/to these chains is also possible and appears to be analogous to the extraction or insertion of lithium in NASICON-type materials. [ 160 , 173 ] Similarly to NASICON, Li + motion is predominantly exercised through 1D wavy pathways along a [010] crystallographic axis [ 174 , 175 ] (this fi nding, which looks natural from the schematic structure image of the olivine-type material, is shown in Figure 18 ); a careful consideration of Li + dynamics in the olivine lattice also reveals that there is Li + hopping between different [010] pathways as well. [ 176 ] However, the diffusion coeffi cient for the displacement along [010] path-ways is two orders of magnitude higher than for the displace-ment between these pathways. [ 177 ]

Tavorite-family poly-anion materials crystallize in triclinic (tavorite, triplite, maxwellite) and orthorhombic (sillimanite) structures; the family is named after the mineral tavorite (LiFe-PO 4 OH). Comprising isostructural lithium fl uoro(hydroxyl) phosphates/sulfates with a variable F − /OH − content ratio, the

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tavorite structure constitutes linear chains of corner-sharing MO 4 F 2 octahedra propagating along the C axis interconnected by corner-sharing SO 4 or PO 4 tetrahedra along the a and b axes, and this design forms a 3D network of Li-conducting tun-nels. [ 178 , 179 ] Other materials of the tavorite-family are variations of this basic 3D framework.

By now, the lithium transport mechanism is under dispute concerning the specifi c mobility along different space direc-tions in the tavorite-type lattice. Calculations [ 180 ] demonstrate that hopping activation energies (and thus, the mobility) along all three [100], [010], and [111] directions are quite similar in the case of LiFeSO 4 F (resulting in 3D lithium conductivity), whereas the work by the Ceder group [ 161 ] predicts that this energy is substantially lower along the [111] axis than along other axes, which results in 1D conductivity. Apparently, as far as the implementation in Li-ion cells with a long cycle life is concerned, a complete substitution of fl uorine for hydroxyl is advantageous since it avoids the undesirable irreversible elec-trochemical reaction of the hydroxyl group with Li.

Considering their electronic structure, poly-anion cathode materials are a promising class of compounds in terms of high operational voltage. Strong covalent bonding within the poly-anion reduces the covalent bonding to the iron ion (induc-tive effect), [ 167 , 181 , 182 ] which, in turn, lowers the redox energies of M + (n) / M + (n + 1) and

.M ( + n) /

.M + (n + 1) . The higher the covalent

component of the bonding within the poly-anion, the lower the position of the metal cation redox level is against the Li/Li + energy level. [ 158 ] The average discharge voltage is 2.98 V for LiFeO 2 cathode, [ 183 ] whereas the discharge voltage is 3.5 V for LiFePO 4 cathodes; [ 184 ] the same effect is demonstrated by the cobalt-based redox couple: the discharge voltage is ∼ 4.1 V for a LiCoO 2 cathode but it is 4.8 V for a LiCoPO 4 cathode. [ 185 ]

This inductive effect assumes that intercalation/de-intercala-tion voltages may be tuned by changing poly-anions. Particu-larly, ab-initio calculations suggest that in the case of materials such as LiCoXO 4 (X = P; As; Sb), Li 2 VOXO 4 (X = P; As; Si; Ge) and Li 2 MXO 4 (M = Mn, Fe, Co, Ni and X = P, Si, Ge) the intercalation/de-intercalation voltages change in the following order: VGe O4 < VSi O4 < VSbO4 < VAs O4 < VP O4 ; generally, this sequence is confi rmed experimentally. This order is observed for vanadium-based materials, [ 186 ] and in the case of iron–based materials V dis char ge

Si O4 ∼ 2.75 V [ 187 ] and V dis char ge

P O4 ∼ 3.5 V, [ 184 ] while

on cobalt, experimental results show that V dis char geSi O4

is between 4.1 V [ 188 ] and 4.25 V, [ 189 ] V dis char ge

As O4 is ∼ 4.6 V, [ 190 ] and V dis char ge

P O4 is

∼ 4.8 V. [ 185 ] Considering poly-anion cathode materials, it is well to bear

in mind that poly-anions bring some inactive mass, which com-promises the specifi c charge capacity. Atomic weights of Si (at. wt. 28), S (at. wt. 32) and P (at. wt. 32) are close enough and all these atoms are substantially lighter than Ge (at. wt. 73) and As (at. wt. 75). Thus, the implementation of the PO 4 anion looks favorable from the point of view of high voltage as well as high specifi c capacity. Since boron (at. wt. 11) is substantially lighter than all the above elements, and the structure of LiMBO 3 (M = Fe, Co, Mn) offers easy paths for Li + (the structure is built by MO 5 trigonal bipyramids, and BO 3 planes form the 3D M(BO 3 ) open framework), boron-based poly-anion compounds have attracted the attention of theoreticians [ 191 , 192 ] and experimen-talists. Unfortunately calculations have revealed that values of

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Figure 19 . LiMnPO 4 (black line, data acquired from) [ 195 ] and LiMnBO 3 (red line, data acquired and replotted) [ 194 ] - cathodes, C/20 rate; the data were acquired using Grafula curve tracking coder). [ 196 ]

1.5

2

2.5

3

3.5

4

4.5

0 10 20 30 40 50 60 70 80 90 100

Volt

age

vs. L

i/Li

+ (V

)

mAh g-1

LiMn(PO4) dashed+triangles, LiMn(BO3) solid+squares, C/20

LiMn(PO4)

LiMn(BO3)

intercalation voltages for boron compounds are substantially lower (0.4 V) [ 192 ] than for phosphorus compounds, and experi-mental results confi rm these fi ndings ( Figure 19 ). [ 193–195 ]

An alternative method to tuning the covalent nature of the metal–anion bonding is to introduce fl uorine into the anion; indeed, if an oxygen ion is replaced by fl uorine, the covalent metal–fl uorine bond would have more ionic character, stabi-lizing the energy of the anti-bonding d-orbital of the metal ion and, thus, the lithium insertion voltage would increase. [ 197 ] The introduction of fl uorine may result in a variation of the poly-anionic material crystal structure (compared with the non-fl uorinated structure), which, in turn, may affect the value of the intercalation/de-intercalation voltage. For example, NASICON-type structured Li x Fe 2 (SO 4 ) 3 (both monoclinic and rhombo-hedral modifi cations) demonstrates V discharge = 3.6 V; [ 159 ] this value is close to V discharge of the tavorite-type LiFeSO 4 F material (3.55 V [ 198 ] and 3.6 V) [ 199 , 200 ] but the same material with a triplite-type structure demonstrates V discharge = 3.9 V. [ 162 ]

When applying ab-initio calculations, it worth remem-bering that they are very useful for screening out promising candidates for materials with operational voltages in the 5 V vicinity, [ 161 , 182 , 201 ] but they might be at variance with experi-mental data on cathode compounds. For example, the calcu-lated value of the FePO 4 F → LiFePO 4 F intercalation voltage is over 5 V [ 161 ] but the experimental study did not confi rm an elec-trochemical activity of this compound above 4 V). [ 199 ]

3.6.1. Co-Based Poly-Anion High Voltage Cathode Materials

The fi rst candidates for high-voltage electrodes are poly-anion materials containing cobalt. Cobalt-based olivine LiCoPO 4 is reported to have a discharge plateau at 4.8 V [ 185 , 202 , 203 ] and LiCoAsO 4 -olivine compound is reported as having discharge voltage expectedly lower, close to 4.6 V. [ 190 ] The addition of

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a small amount of vanadium margin-ally enhances the average voltage of the Li 1 + 0.5x Co 1 − x V x (PO 4 ) 1 + 0.5x electrode (x = 0.05 and 0.1) [ 204 ] and the addition of a small amount of iron stabilizes the crystal struc-ture [ 205 , 206 ] and improves Li + transport in the material. [ 207 ] This enhances the cycling and charge capacity of LiCoPO 4 without compromising the discharge voltage. Cobalt-containing olivine LiMn 1/3 Fe 1/3 Co 1/3 PO 4 is reported as having a plateau at 4.9 V [ 208 ] (the average operational voltage of the compound is reported to be about 3.72 V, [ 209 ] though). As the iron content (y) increases, the oliv-ines LiFe y Co 1− y PO 4 demonstrate shrinkage of the Co + 2/ + 3 -related plateau, whereas the olivines LiMn 0.25 Co 0.75 PO 4 nearly dem-onstrate solely (Co + 2/ + 3 )-related plateau (the position of which shifts a bit, from 4.83 V to 4.73 V compared with LiCoPO 4 material), albeit with an extended charge capacity. [ 210 ] A marginal amount of Mn (but not Ni or Mg!), y ≤ 0.1, enhances the cycling capability of LiMn y Co 1− y PO 4 ; [ 211 ] LiCoPO 4 F has a voltage plateau even closer to 5 V then

LiCoPO 4 [ 212 ] and the theoretical estimation indicates the interca-lation voltage of ∼ 4.9 V for cathode with -SO 4 F anion. [ 161 , 213 ] At the same time, Co-based cathodes with -SiO 4 , [ 188 , 189 , 214 ] -BO 3 , [ 194 ] and -V 3 O 8 [ 215 ] anions were found to have distinctly lower voltages (all these results agree well with ab-initio calculations). [ 191 , 216 ]

3.6.2. Ni-Based Poly-Anion High Voltage Cathode Materials

Calculations predict that Ni-containing poly-anion materials may have the highest intercalation voltages, over 5 V. [ 161 ] At the same time, the experimental results are ambiguous. The calculated intercalation voltages for LiNiSO 4 F are 5.16 V [ 217 ] and 5.35 V, [ 161 ] whereas the electrochemical activity of nickel in this compound was not confi rmed experimentally. [ 199 ] The discharge voltage of Li in LiNiPO 4 was estimated at 5 V, [ 22 , 218 ] but its electrochemical activity was not confi rmed, [ 219 ] and later the same authors found some indications of Li + intercalation into LiNiPO 4 in the 5 V-region. [ 220 ] Experimentally, LiNiPO 4 F discharge voltage is demonstrated to be close to 5.3 V, [ 221 ] and ab-initio calculations predict a value of 5.5 V. [ 161 ] The calcu-lated discharge voltages for oxy-compounds LiNiO(PO 4 ) and LiNiO(SO 4 ) are 5.27 and 5.01 V, correspondingly, [ 161 ] but we are not aware of experimental data on the synthesis or electro-chemical properties of these compounds.

3.6.3. Low Conductivity: A Problematic Point for Poly-Anion Cathode Materials

The distinctive (and problematic!) feature of polyanion cathode materials is their poor electronic and ionic conductivity. [ 62 , 222 ] Indeed, typical room temperature electronic conductivity and Li + diffusivity (ionic conductivity is proportional to ionic dif-fusivity) [ 223 ] of layered oxide cathode materials and spinel-type compounds are in the range of ρLiCo O2 ∼ 10 − 3 S cm − 1 and

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DLiCo O2 ∼ 10 − 9 cm 2 s − 1 , [ 88 , 224 ] ρLi Mn2 O4 ∼ 10 − 4 S cm − 1[ 225 , 226 ] and DLi Mn2 O4 ∼ 10 − 9 cm 2 s − 1 , [ 227 , 228 ] whereas these values for poly-anionic compounds are several orders of magnitude lower; they are in the range ρLi F e P O4 ∼ 10 − 8 –10 − 11 S cm − 1 and DLi F e P O4

∼ 10 − 11 –10 − 17 cm 2 s − 1 . [ 177 , 229–232 ] It is commonly suggested that the limiting factor in th epoly-anion materials’ electrochemical per-formance is mostly the ionic and/or electronic transport within the particles; low diffusivity and conductivity limit material cur-rent and charge performance, decreasing the high values of intercalation/de-intercalation voltages. [ 233 , 234 ]

Several strategies have been employed to mitigate the problem. One approach is to enhance the specifi c conductivity of the compound. [ 235 ] It was demonstrated that a small polaron transport is responsible for the electronic conductivity, [ 231 , 236 , 237 ] and by introducing supervalent dopants (such as Mg 2 + , Al 3 + , Ti 4 + , Zr 4 + , Nb 5 + , W 6 + , etc.), which substitute for Li + , it is pos-sible to produce an excess of free, unbound lithium vacancies in the poly-anion cathode material (the bulk of these studies deal with LiFePO 4 ); [ 234 ] since the neutrality would require a Li defi ciency in the material matrix, the substitution introduces additional charge carriers. By this strategy, room temperature electronic specifi c conductivity can be increased [ 238–241 ] up to a factor of ∼ 10 8 as compared to the starting material, reaching values of ∼ 10 − 2 S cm − 1 . [ 238 ] At the same time, there are some arguments challenging the electronic effect of such dopants (located primarily on the Li sites) and favoring the conclusion that these dopants hinder Li diffusion. [ 239 ] Regarding electro-chemical properties, up to now, LiFePO 4 doping has not been revealed to hold signifi cant advantages over the undoped materials. [ 235 ]

The other strategy is to diminish material grain size, since it is the common concept that the smaller the grain is, the lower the overall grain resistance. [ 184 , 233 , 243–245 ] The limiting factor of poly-anion cathode performance is, though, low ionic diffusivity. Typically, the ionic conductivity of poly-anion material is several orders of magnitude lower than the electronic conductivity (in any case, larger than 10 − 11 S cm − 1 and lower then 10 − 12 S cm − 1 , [ 246–248 ] correspondingly). Based on this fact, it was suggested that the role of carbon coating in facilitating electron transport is to ensure that each grain contacts an adequate number of carbon black particles. Thus, carbon coating material particles has little or no effect on charge transfer. [ 243 , 249 ]

However, such important parameters as particle wiring quality and total cathode resistance depend on the amount and method of conductive additive (carbon) introduction. [ 250 , 251 ] This circumstance justifi es the development of this fi eld and explains the implied performance improvements, [ 252 , 253 ] such as rate capability improvement [ 254 , 255 ] and cycle life improve-ment. [ 256 ] Particularly useful is carbon coating, which helps to cure the most problematic features of nanoparticle electrodes, namely poor particle contact and high reactivity at the surface, retaining the advantages of nanoscale technology. [ 244 ] Curing surface layer disorder allows the improvement of Li + transport through this layer and improves contact (intergranular) mate-rial conductivity, while the poor contact conductivity may cause not only low current collection but also intergranular contact, overheating and, thus, material decomposition. [ 257 ]

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3.7. High Voltage Capabilities: Spinel-Type vs. Poly-Anion Oxides

Up to now, the most promising development in nickel-doped Mn spinel LiNi 0.5 Mn 1.5 O 4 allows an operational voltage in the range of 4.7–4.8 V [ 111 ] and also presents a high cyclability (the material has demonstrated that only a marginal capacity fade after 500 cycles at 10C rate being modifi ed with a trace amount of ruthenium). [ 258 ] High voltage Co-based poly-anion oxides show substantially lower practical charge capacity than spinels [ 149 , 195 ] but exhibit a discharge activity at voltages close or even over 5 V. [ 212 ] At the same time, high-voltage Co-based poly-anion materials demonstrate substantial performance fade just after 30 cycles. [ 190 , 195 ] Scant information is available on high voltage Ni-based poly-anions; it appears that these mate-rials have a heavy cycle fade. [ 221 ] Regarding safety, high-voltage spinel type oxide LiNi 0.5 Mn 1.5 O 4 demonstrates a fair thermal stability, up to 250 ° C, [ 150 ] whereas the thermal decomposition of LiCoPO 4 starts at temperatures below 200 ° C. [ 259 ] Data on the thermal stability of Ni-based poly-anion compounds are not available.

The bottom line is that, although poly-anion compounds offer considerable potential for extending the cathode working area up to 5 V (vs. Li/Li + ) and even beyond that, currently these compounds are exceeded in practical charge capacity, cyclability and safety by spinel-type high voltage LiNi 0.5 Mn 1.5 O 4 –based cathode materials. Much work is to be conducted on the refi ne-ment and enhancement of these parameters of Co- and Ni-based high voltage cathode materials.

4. Conclusion

The development of 5 V cathode materials is currently focused on meeting the cycling, current and charge capacity challenges. Oxide spinels and poly-anion materials with Ni and Co cations are currently considered the most promising for future energy storage requirements. Until now, the most successful high-voltage cathodes are materials based on the LiNi 0.5 Mn 1.5 O 4 spinel. Full utilization of the potential of these materials in terms of cyclability, current performance and safety continues to be a challenge. Primary emphasis is now being placed on the improvement of their structural stability by introducing dopants and also on improvement of charge currier transport by sizing the material grains down to the nanometer scale. The investigation of the infl uence of anion alteration is also underway. Substantial progress has been made in the fi eld of cathode composition, microstructure and morphology, however, additional efforts are indeed needed.

Acknowledgements Support for this work was provided by the Grand Technion Energy Program (GTEP), EU FP InnoEnergy Project, and the Leona & Harry B. Helmsley Charitable Trust.

Received: January 26, 2012 Published online: June 18, 2012

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