Effects of heteroatoms on electrochemical performance of electrode materials for lithium ion batteries

Electrochimica Acta 47 (2002) 3491
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Effects of heteroatoms on electrochemical performance of electrode materials for lithium ion batteries Y.P. Wu a,b,*, Elke Rahm b, Rudolf Holze b,1 a

b

Division of Chemical Engineering, INET, Tsinghua University, Beijing 102201, China Chemnitz University of Technology, Institute of Chemistry, AG Elektrochemie, 09107 Chemnitz, Germany Received 12 February 2002; received in revised form 14 May 2002

Abstract Recent studies of lithium ion batteries focus on improving electrochemical performance of electrode materials and/or lowering cost. Doping of active materials with heteroatoms is one promising method. This paper reviews the effects of heteroatoms on anode materials such as carbon- and tin-based materials, and cathode materials such as LiCoO2, LiNiO2, LiMn2O4 and V2O5. There are favorable and unfavorable effects, which depend on the species and physicochemical states of heteroatoms and the parent electrode materials. In the application of lithium ion batteries advantageous factors should be exploited, unwelcome side effects should be avoided as far as possible. Considerable gains towards improved electrochemical performance of the electrode materials have been achieved. Nevertheless, there are still problems needing further investigation including theoretical aspects, which will in the meanwhile stimulate the investigation for better electrode materials. # 2002 Elsevier Science Ltd. All rights reserved. Keywords: Lithium ion batteries; Electrode materials; Heteroatoms; Electrochemical performance

1. Introduction The beginning of research on lithium secondary batteries goes back to the 1960s and 1970s due to the first energy crises and the growing interest in power sources for mobile applications. However, no breakthrough has been made before the early 1990s though there were a lot of demands for high power and energy density battery systems. The main problem is the safety associated with the formation of dendrites during repeated charge and discharge. In 1990, Sony Corporation achieved great progress on this safety problem through substituting the lithium metal electrode by carbonaceous material, and introduced the concept of the ‘lithium ion battery’. Since this time, its development has been very rapid due to its many advantages over traditional rechargeable battery systems such as average high output voltage (up to 3.7 V), light weight, high energy density, excellent * Corresponding author. Fax: /49-371-531-1371 E-mail addresses: [email protected] [email protected] (R. Holze). 1 ISE member.

(Y.P.

Wu),

cycling life (above 1000 cycles), low self-discharge (lower than 5% per month) and the absence of potentially environmental pollutants such as lead and cadmium [1]. Now they are used as power sources in many kinds of high value electronics such as video cameras, portable computers and telephones, and its application such as in zero-emission vehicles, medical instruments, aerospace industry and military is almost reality. Its market is very promising, but the competition is intense. As a result, electrode materials have been widely investigated, and very rapid and significant improvements have been made over the electrode materials prepared in the early 1990s including short manufacture time, lower preparation temperature, lower cost, high capacity, good cycling behavior [1]. These kinds of improvements are due to the employment of novel preparation technology such as sol
/gel or template approaches, and introduction of heteroatoms. However, some methods are only limited to special electrode materials. One widely adopted promising method is the introduction of heteroatoms into electrode materials. The effects of the introduced heteroatoms on electrochemical performance are complicated depending on species and forms of heteroatoms, and parent structures. Here we review recent progress on

0013-4686/02/$ - see front matter # 2002 Elsevier Science Ltd. All rights reserved. PII: S 0 0 1 3 - 4 6 8 6 ( 0 2 ) 0 0 3 1 7 - 1

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this aspect. The electrode materials for anode and cathode are treated separately.

2. Anode There are several kinds of anode materials: carbonbased, tin-based and others. 2.1. Carbon-based anode materials As mentioned above, carbonaceous materials initially used as the anode for lithium ion battery were disordered non-graphitized carbons (hard carbon). Later graphitic carbon was also found to be able to act as good anode material. Since these materials show good reversibility for lithium intercalation and de-intercalation and are non-toxic and low cost, they are now widely employed [1]. In the middle of the 1990s, it was found that amorphous carbon prepared at low temperatures could be a promising anode material for the secondgeneration of lithium ion battery since its preparation temperature is lower and its capacity higher than that for graphitic carbon [1,2] though it has not yet come to commercialization so far to our knowledge. A lot of dopant elements, which can be grouped in two categories: nonmetal and metal elements, have been introduced into graphitic and amorphous carbons. 2.1.1. Introduction of nonmetal elements The effects of nonmetal elements are reviewed following the order in the periodic table. Hydrogen is the lightest element, and there are different opinions concerning its effect on electrochemical performance of amorphous carbon materials. Some researchers observed that there was a linear relationship between the ratio of H/C and reversible capacity [3] and suggested that hydrogen was involved in interaction with lithium during discharge and charge processes [4] according to: C
H2Li 0 C
LiLiH 2C
H2Li 0 2C
Li1=2H2

(1) (2)

However, the content of H can be as low as 0.039, and the reversible capacity can be up to 975 mA h g1 [5], which is difficult to explain assuming a contribution of hydrogen. In addition, with some kinds of carbonaceous materials, it is found that the content of hydrogen in the carbon structure does not change after several cycles [6]. Evidently, the action of hydrogen is still not clear. Perhaps the linear relationship is just a coincidence since hydrogen atoms cannot be completely removed during heat-treatment of organic precursors below 1200 8C. At most, it is active only with specific carbons. Boron is the only nonmetallic element in Group IIIA. It can be incorporated into carbon materials in two different forms: in atomic form [7
/9] or in the form of

compounds [8
/10]. The first one is the result of pyrolysis of boranes or other boranic compounds with hydrocarbons. Carbon and boron atoms can be deposited together when preparing the carbonaceous materials by chemical vapor deposition (CVD) [7]. The latter form is obtained by adding directly boronic compounds such as B2O3 and H3BO3 into precursors followed by heat-treatment of the mixtures [10
/12]. The enhancement of the reversible lithium capacity after introduction of boron atoms seems to be due to the electrondeficiency of boron [7,10
/12]. It is an electron acceptor and can increase the combination energy of lithium with carbonaceous materials from E0 to E0/D (E0 is the combination energy for lithium intercalation into graphite resulting in Lix C6). In addition, it is favorable for the graphitization process, lowers the number of edge planes with dislocations [13], and decreases the interlayer distance d002 [14]. Its effect on charge voltage (deintercalation) is situated at about 1.1
/1.6 V versus Li / Li (in the following text potentials are referred to Li /Li if not indicated otherwise). Of course, its influence on electrochemical performance of the prepared carbonaceous materials depends on various experimental conditions such as the heat-treatment temperature. If boron exists in the form of boron carbide the irreversible capacity will increase [8,9]. The effects of both forms on reversible lithium capacity of the carbonaceous materials are slightly different. In the first form, reversible capacity increases linearly with the content of boron up to 9%, and the boron content can be up to 13% [7]. With the second form, the maximum reversible capacity is obtained with about 1.0
/2.0% added boron, and the irreversible capacity decreases after the introduction of boron [10
/12]. Effects of boron were also calculated based on a semiempirical molecular-orbital model [15,16]. Results show that the introduction of boron is favorable for lithium intercalation. When a layer of BCx is coated onto the surface of natural graphite, the performance improves considerably [17]. In contradiction, another theoretical study based also on a semi-empirical molecular orbital method concludes that the substitution of the carbon by boron is not effective for lithium storage [18]. This illustrates the complexity of carbon structure. These results suggest that the bonding states of boron markedly influence the properties of the carbonaceous materials. Nitrogen can also be incorporated into carbonaceous materials. However, there are contradicting reports [19
/ 21]. It was first reported that nitrogen existed in carbonaceous materials in two forms, which were called chemical nitrogen and lattice nitrogen [19]. The former can irreversibly react with lithium and causes an enhancement in irreversible capacity. As a result, carbonaceous materials doped with nitrogen were not regarded as suitable anodes for lithium secondary

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Table 1 Relationship between the reversible capacity and the nitrogen content in carbonaceous materials Precursors

N/C atomic ratio

Reversible capacity (mA h g 1)

References

Benzene Pyridine Pyridinechlorine Pyridinechlorine Polystyrene Poly(4-vinyl pyridine) Polyacrylonitrile Melamine resin

0 0.0800 0.0855 0.137 0 0.0804 0.195 0.217

249 335 392 507 345 386 418 536

[20] [20] [20] [20] [21] [21] [21] [21]

batteries. However, quite different results using the same CVD process and nitrogen source (pyridine) were achieved elsewhere as shown in Table 1 [20]. The reversible capacity increases with the nitrogen content. In pyrolytic carbons obtained from nitrogen-containing polymers, reversible capacity also increases with the nitrogen content and can be above the theoretical value of graphite 372 mA h g1 [21]. According to results of X-ray photoelectron spectroscopy nitrogen also exists in two forms in carbonaceous material, which are different from those reported in [19]. The two forms are graphene nitrogen and conjugated nitrogen. The former is located in graphene molecules, and its N1s binding energy peak is at 398.5 eV; the latter is in the form of
/C /N
/, which is not incorporated into the graphene molecules, and its N1s binding energy peak is at 400.2 eV [21]. One example is shown in Fig. 1 [21]. According to X-ray absorption spectra the chemical nitrogen is perhaps in the state of an amine-group [19]. It is known that amine-group nitrogen is very active and can result in irreversible capacity. Later on a wide range of nitrogen-containing polymers has been investigated as precursors to prepare anode materials [22,23]. Effects of nitrogen were also calculated based on semi-empirical molecular orbital models [15,16,18], the results could not definitely deny the favorable effects of

nitrogen [15,16]. Obviously there are limits in computational chemistry since the carbon structures are very complicated. Modelling has not yet arrived at a level completely reflecting especially the binding forms of heteroatoms, various states and microstructures of carbons. Consequently, these kinds of computations are considerably speculative. Recent result evidently showed this limit again [8,9]. This also shows that the bonding state of the heteroatoms markedly influences the properties of the carbonaceous materials. When a layer of Cx N is coated on the surface of graphite, capacity and cycling behavior of the composite improve [24], which also shows the favorable effects of the doped nitrogen. The effects of oxygen (Group VIA of the periodic system) and fluorine (Group VIIA) doping into the carbon structure has not been studied. However, they can be attached to the surface of carbons, and their effects on electrochemical performance can be ascribed to the modification of the surface structure of carbons such as elimination of imperfections, formation of a dense layer of oxides/fluorides and production of micropores/nanochannels [4,25
/29]. Consequently, the cycling behavior and reversible capacity strikingly improve (Fig. 2). Silicon can also be introduced into carbonaceous materials, both via CVD methods [30] and pyrolysis of

Fig. 1. X-ray photoelectron spectra of both kinds of nitrogen N1s in carbon from heat-treatment of poly(acrylonitrile) at 600 8C (from reference [21]).

Fig. 2. Cycling behavior of natural graphite A unmodified and modified by oxidative aqueous solutions of H2O2, Ce(SO4)2, HNO3 and (NH4)2S2O8 (from references [28] and [29]).

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silicon-containing polymers such as poly(methylphenyl silane) and their mixtures with other polymers [31
/34]. Both methods can be performed at low temperatures (B/1200 8C), and the obtained carbons are amorphous. The distribution of silicon is on a nanometer scale. When the introduced amount of silicon is within the range of 0
/6% (molar), reversible capacity increases from 300 (i.e. that of pure carbon) to 500 mA h g1, and the increase in reversible capacity is about 30 mA h g1/ 1% of silicon, which indicates that each introduced silicon atom can reversibly react with 1.5 Li atoms. Its effect on capacity is observed at electrode potentials between 0.1 and 0.6 V. Reversible capacity does not fade fast during cycling. The maximum amount of silicon can be 30% [34]. Recent research work from Sony Corporation showed that an amorphous carbon prepared by heat-treatment of bamboo at low temperature can have a reversible capacity as high as 600 mA h g1 [35], and the main reason is the existence of silicon according to X-ray diffraction measurements. However, later investigation revealed that silicon exists not in the form of nanodispersed silicon, but in the form of nanodispersed silicon oxycarbide clusters [36] though the results of electrochemical performance did not change. As a result, stoichiometric compositions of Si, C and O in a tri-phase diagram should be avoided for applications in lithium ion batteries as anode materials [33]. The mechanical mixing of silicon and graphite produced a composite of C0.8Si0.2 whose reversible capacity can be up to 1039 mA h g1 and after 20 cycles can be still 794 mA h g1 [37]. A composite of silicon and carbon can also yield considerably increased reversible capacity [38]. The main reason seems to be that the introduced silicon can promote the diffusion of lithium in the interior of carbonaceous materials and effectively prevent the production of dendrites. The increase of reversible capacity cannot be completely ascribed to the formation of Lix Si alloy because silicon is present in part in the form of Si
/O compounds [36], and further investigation is required. The effects of phosphor on carbonaceous materials depend on the precursors [39
/43]. The introduction of phosphor in pitch coke mainly affects the surface structure of the carbonaceous materials [39], phosphor interacts preferably with edge planes of the carbonaceous materials. Since its atomic diameter (310 pm) is larger than that of carbon (154 pm), this kind of interaction results in an increase of the interlayer distance d002, and thus lithium intercalation and deintercalation are favored. When phosphoric acid reacts initially with precursor materials, chemical bonds may form involving surface functional groups, and subsequent heat-treatment of the homogeneous mixture can result in complete incorporation of phosphor into the carbonaceous structure [40,42]. Results from X-ray photoelectron spectroscopy illustrate that phosphor

exists in a single form, which bonds not only with carbonaceous material but also with oxygen atoms due to low heat-treatment temperature (B/1200 8C). The introduced phosphor affects both the electronic state and the carbonaceous structure, this also depends on the precursors and heat-treatment conditions. At relatively high temperatures (]/800 8C), its introduction is favorable for an increase in reversible capacity. When phosphor is introduced into epoxy resin, it softens the carbon structure and favors the graphitization process to produce larger graphite crystal. As a result, the reversible capacity is up to 550 mA h g1 and the coulombic efficiency in the first cycle is 83% [43]. The effects of the incorporation of sulfur on the electrochemical performance of the obtained carbons depend on the sulfur states therein [44
/47]. X-ray photoelectron spectroscopy illustrates that S atoms can exist in three states:
/C
/S
/C
/,
/C
/S
/S
/C
/, and sulfate/sulfonate, and the binding energy peaks of S2p are situated around 164.1, 165.3 and 168.4 eV, respectively. The introduced sulfur affects also the carbon structure. Table 2 indicates that reversible capacity after the introduction of sulfur is increased considerably. The charge (de-intercalation) curve shows that the plateau below 0.5 V lengthens [45]. Of course, the effect of incorporated sulfur also depends on the heat-treatment conditions such as type of precursor, temperature and soak time. For example, the addition of sulfuric acid into the precursor PAN prior to heat-treatment introduces favorable and unfavorable factors although in effect both reversible and irreversible capacities increase [44]. 2.1.2. Introduction of metallic elements Metals introduced into carbonaceous materials include main group and transition group elements. Main group elements studied so far include potassium (Group IA), magnesium (IIA), aluminum and gallium (IIIA); transition metals include vanadium, iron, cobalt, nickel and copper. Potassium is introduced into the carbonaceous materials via the formation of the intercalation compound KC8 [48], which is used as electrode material. After the de-intercalation of potassium, lithium intercalates into graphite instead of potassium. Since the interlayer distance d002 after the intercalation of potassium (341 pm) is larger than that of graphite (336 pm), rapid intercalation of Li is favored. The intercalation can proceed up to the level of LiC6, and reversible capacity can be up to 372 mA h g1. In addition, with KC8 used as anode material there are more choices available in cathode materials such as some low cost, Li-free compounds. The introduction of magnesium was found accidentally. When coffee pea is heat-treated, the obtained amorphous carbon can reversibly store lithium up to

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Table 2 Effects of added sulfur on the reversible capacity of the carbon anode for lithium ion secondary batteries Sulfur content in carbons (%)

Reversible capacity below 0.1 V (mA h g 1)

Reversible capacity below 0.5 V (mA h g 1)

Reversible capacity below 3.0 V (mA h g 1)

0 2.41

105 269

245 412

363 568

670 mA h g1. In X-ray diffraction patterns diffraction peaks were assigned to magnesium and its compounds [35]. However, details of its effect are not yet understood. The introduction of aluminum and gallium can increase the reversible capacity of carbon anodes mainly because they can form solid solutions in planar structures with graphene molecules. Their Pz channels are empty, and thus more lithium can be stored [49,50]. Aluminum may be deposited on the surface of graphite [51]. Due to a sharp decrease in charge-transfer resistance during cycling, reversible capacity and rate capability of the coated graphite anode improve markedly (Fig. 3). Transition elements such as vanadium [52
/54], nickel and cobalt [55] are added in the form of oxides into the carbonaceous precursors, and then heat-treated. Since they act as catalysts favoring the graphitization process and increasing the interlayer distance d002, reversible capacity increases and cycling behavior improves. For example, vanadium is added in the form of V2O5. During the heat-treatment process it coordinates with the formed graphene molecules and forms a complex molecule, VO(graphene)2. This kind of complex acts as a nucleation agent to get a more ordered arrangement of graphene molecules. This process is schematically illustrated in Fig. 4. Oxides such as those of nickel, cobalt and iron [56] can be deposited onto the surface of graphite and act as catalysts for the chemical oxidation eliminating active defects. They can also act as matrices for lithium storage

Fig. 3. Cycling performance of natural graphite (curves d, e and f) and Al-treated sample (curves a, b and c): circles, triangles and rectangles represent 0.2, 0.5 and 1 C rate, respectively (from reference [51]).

Fig. 4. Schematic diagram of the production of carbonaceous material by heat-treatment of a mixture of V2O5 and a polymer and the formation of VO(graphene)2 (from reference [54]).

after the first irreversible reaction resulting in Li2O and metals. Metallic nickel can be deposited on the surface of graphite [57], and block part of the edge surface exposed to the electrolyte solution. This minimizes intercalation of solvated lithium at these edge sites. As a result, the decomposition of electrolytes and exfoliation of graphene layers are minimized. Coulombic efficiency in the first cycle increases from 59 to 84% and reversible capacity increases by about 30
/40 mA h g1. Copper and iron can be doped into graphite though this process is complicated [58]. At first natural graphite reacts with the chlorides at high temperature and forms intercalation compounds, which are later reduced with LiAlH4. After these treatments, the interlayer distance increases and the surface of the material with exposed edge sites becomes smoother. As a result, electrochemical properties become better, selected results are shown in Table 3 [58]. When the content of M in these doped compounds Cx M (M /Cu and Fe) is too large (x B/24), intercalation sites for lithium are less numerous and reversible capacity decreases. On the contrary, when x /36 and the content of M is too small, the irreversible capacity in the first cycle enhances and poor overcharging behavior is found. Copper and its oxides can be also coated onto the surface of graphite [59,60]. The coated copper can remove/cover some active sites and decreases the absorption amount of water under high humidity condition (1000 ppm H2O), and thus the composites exhibit excellent cycling behavior when cells

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Table 3 Electrochemical properties of the carbonaceous materials doped with Cu and Fe Carbonaceous materials

Natural graphite C22Cu C24Cu C30Cu C22Fe C24Fe C30Fe

Discharge capacity (mA h g 1)

611 449 463 491 398 412 433

are built under relatively high humidity atmosphere (Fig. 5) [59]. Further metals can form compounds X
/C or Li
/X
/ C (X includes Zn, Ag, Mg, Cd, In, Pb and Sn), which can also be used as anode materials for lithium ion batteries [61]. There is an evident improvement in the electrochemical performance. The main reason is that the introduced metals favor Li diffusion. The above results pertaining to the introduction of heteroatoms into carbonaceous anode materials show that there are two kinds of factors: favorable and detrimental. The favorable effects include: (1) acting as host for lithium storage leading to enhancement of reversible capacity such as Si, S [30,37,45], (2) decreasing electron density of the host structure and thus more lithium can be intercalated such as B, Al, Ga and N [7,10
/12,21,49,50], (3) favoring more ordered arrangement of graphene molecules during heat-treatment such as P and V [42,53] or the graphitization process such as B [13], (4) increasing interlayer distance, d002, such as K, V and P [40,42,48,53], (5) increasing electronic conductivity [51] and/or favoring lithium diffusion [38,61], (6) removing/covering active imperfections [26
/29,59], (7) forming a dense layer at the surface acting as good passivating film [26
/29], (8) introducing of nanochannels and/or micropores [4,26
/29]. Consequently, elec-

Charge capacity in different cycle (mA h g 1) 1

20

50

100

385 395 417 408 398 412 433

343 392 410 406 396 408 430

327 390 407 406 396 403 429

327 390 407 406 396 403 429

trochemical performance including reversible capacity, coulomb efficiency in the first cycle, cycling behavior and rate capability is improved. Unwelcome factors are mainly due to the following reasons: (1) side reactions with lithium such as amine nitrogen [19,44], (2) unfavorable for graphitization such as introduction of P below 600 8C [40]. Of course, other kinds or two or more types of heteroatoms can be introduced. 2.2. Tin-based anode materials Research on tin-based anode materials was first initiated by Fuji Film Company in the middle 1990s [1], and later moved from oxides to alloys. Their advantages are easy preparation such as heat-treatment below 1500 8C or simple mechanical milling and high volumetric capacity, which can be twice of that of carbonaceous anode materials. As far as we know, these materials are still waiting for commercialization. However, understanding the effects of heteroatoms will provide some valuable clues for further investigation and perhaps breakthrough will come in the near future. 2.2.1. Oxides Tin oxide and stannic oxide can be matrices for lithium storage since tin can form alloys with Li up to Li22Sn5. However, Li reacts first with the oxides forming Li2O and then forms alloy with the produced tin according to equations 3 and 4, [1]: 6 LiSnO2 =SnO 0 2Sn3Li2 O Snx LiULix Sn (x5 4:4)

Fig. 5. Cycling behavior of natural graphite and graphite with deposited copper when assembled into test cells under high humidity condition (about 1000 ppm H2O) (from reference [59]).

(3) (4)

Irreversible capacity in the first cycle is too large, and they are not suitable for application in lithium ion batteries. It is reported that the addition of other elements such as B, Al, P, Si, Ge, Ti, Mn and Fe can prepare amorphous composite oxides [62
/65]. This amorphous structure consisting of electrochemically active centers of Sn
/O and the surrounding amorphous network is stable during charge and discharge [63]. The added other oxides in the amorphous network delocalize the Sn(II) active centers and enable effective storage of

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lithium. Reversible capacity depends on the amount of Sn
/O active centers, the maximum can be beyond 600 mA h g1. In addition, the amorphous glass formed by the other added oxides enhances lithium diffusion in comparison with crystalline tin and stannic oxides, thus reversible intercalation and de-intercalation are favored. The density of the composites is higher than that of graphite (3.7 g cm3). Each unit can store eight Li, and its volumetric energy density is above 2200 mA h cm 3, double that of carbon-based anode materials which for amorphous carbon and graphite is less than 1200 and 500 mA h cm 3, respectively [63]. The mechanism for lithium storage is as follows [62,63]: SnB0:5 P0:5 O3 7
8 LiU Li7
8 SnB0:5 P0:5 O3

(5)

Lithium exists in the form of lithium ion in the composite. However, other researchers found that lithium first reacts with the composites to form Li2O, and then forms an alloy with the reduced tin, as shown in equations 6 and 7, [62,64]. As to its specific actions, further studies are still underway. 4 LiSn2 BPO6 0 2 Li2 O2 Sn1=2 B2 O3 1=2 P2 O5 8:8 Li2 Li2 O2 Sn1=2 B2 O3 1=2 P2 O5

(6)

U2 Li4:4 Sn1=2 B2 O3 1=2 P2 O5 2 Li2 O

(7)

2.2.2. Tin-based alloys Metallic tin can form alloys with Li up to Li22Sn5, its reversible capacity is very high. However, there is a drastic expansion and contraction during charge and discharge (sharp change), and the formed alloy Lix M is very brittle, resulting in pulverization of the alloy and fast fading of capacity. When other kinds of inert or inactive metals M? such as Cd, Ni, Mo, Fe and Cu [66
/ 78] are added and alloys like Cu6Sn5d (d /09/1) are formed, lithium can reversibly intercalate and de-intercalate. This is a promising alloy anode for lithium ion batteries [76]. The main reason is that addition of M? results in ductile alloys with greatly reduced volume change [79,80]. A typical example is Cu6Sn5, which has been well studied. When lithium intercalates into Cu6Sn5, there is a phase change proceeding in two steps. First, Li2CuSn is formed. The voltage plateau below 0.4 V indicates the coexistence of Cu6Sn5 and Li2CuSn. When further intercalation proceeds and the voltage is below 0.1 V, a Li-rich phase is produced, Li4.4Sn. In this phase, the Li-rich phase Li4.4Sn and Cu coexist. The de-intercalation also goes via multi-steps resulting in Cu6Sn5. At first Li de-intercalates from Li4.4Sn to produce Li4.4x Sn. With progressing de-intercalation, Li4.4x Sn Sn reacts with copper producing Li2CuSn. Later lithium de-intercalates from Li2CuSn forming Li2x CuSn that

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contains vacancies. With x/1 in Li2x CuSn, deintercalation of lithium proceeds to the initial intermetallic compound Cu6Sn5. The voltage plateau at 0.8 V corresponds to the coexistence of Li2x CuSn and Cu6Sn5. Diffusion of Li in Lix Cu6Sn5 is slow. Its initial gravimetric capacity is 200 mA h g1, smaller than that of graphite, but volumetric capacity is high. For example, its density is 8.28 g cm 3, 200 mA h g1 corresponds to 1656 mA h cm 3 whereas the value for LiC6 is only 850 mA h cm 3. An increase in the content of copper in Cux Sn alloys provides better cycling behavior at the expense of a lower capacity [78]. Ni3Sn2 and Co3Sn2 are similar to h-Cu6Sn5 in structure, and reversible capacity can be 327 or 2740 mA h cm 3 [67], much higher than that of currently available carbonaceous materials. When the utilized capacity is only 1500 mA h cm 3, fading in capacity becomes slow. In the case of the alloy NiSn, reversible capacity is only 77 mA h g1, which is perhaps associated with a low diffusion coefficient of Li [70]. Antimony is added into tin forming multi-phase structures. When the particle size of the alloys is small, the cycling behavior is improved [81
/83]. For example, when it is below 300 nm, capacity can be up to 360 mA h g1 after 200 cycles [81]. During the first lithium insertion, these submicron particles are expanded into a porous material without formation of major cracks [83]. Other kinds of heteroatoms can also be introduced, they should suppress volume change during lithium intercalation and should provide electronic conductivity. Otherwise improvements will be limited.

2.3. Other kinds of anode materials Silicon can also react with lithium to form Lix Si up to a maximum of Li22Si5, whose theoretic capacity is 4000 mA h g1 [84]. However, there is also a drastic change in volume during charge and discharge processes, which leads to a pulverization and results in a fast fading of capacity. When it is dispersed in a solid, conducting and ductile metallic matrix such as chromium [85] or silver [86], the cycling behavior is improved. For example, the alloy SiAg exhibits a reversible capacity of about 280 or 1150 mA h cm 3 and little fading in capacity within the first 50 cycles as shown in Fig. 6. When silicon is doped with chromium to form a composite, reversible capacity enhances from 550 of Li
/Si to 800 mA h g1. Capacity is associated with the initial ratio of Li/Si. When the ratio of Li/Si is about 1:3.5, capacity is at the highest level and good in cycling [85]. TiN can be doped into silicon forming a nanoscale composite. TiN acts as an inert matrix providing structural stability. Though its reversible capacity is low, about 300 mA h g1, it is good in capacity

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LiCoO2. When aluminum is used partly as a substitute for cobalt, the formed LiAly Co1y O2 solid solution has an a-NaFeO2 structure, and can be kept as a single phase up to y /0.5. The open-voltage and working voltage increase with aluminum content. Initial reversible capacity of LiAl0.15Co0.85O2 can be up to 160 mA h g1 and the volume structure does not change after ten cycles [94,95]. Though capacity fading at room temperature is significant, both reversible capacity and cycling behavior at 55 8C are much improved. The observed improvement caused by the doping is from the stabilization of its layered structure. 3.2. LiNiO2 Fig. 6. Charge (lithium de-intercalation) capacity as a function of cycle number for SiAg obtained after milling for 2, 15 and 50 h each. The cells were cycled between 0.07 and 2.0 V. (from reference [86]).

retention [87]. In addition, it can be prepared easily, just by mechanical milling. Behavior of lithium storage in germanium-based alloys is in principle similar to that in tin and silicon [88], which can also form alloys with lithium up to the level of Li4.4Ge: 4:4 LiGeM ULi4:4 GeM

(M : metal dopant)

(8)

and lithium can intercalate reversibly.

3. Cathode materials There are many kinds of cathode materials [1]. Here we focus on four kinds of cathode materials, viz. lithium cobalt oxides (LiCoO2), lithium nickel oxides (LiNiO2), spinel-type lithium manganese oxides (LiMn2O4) and vanadium oxides (V2O5). 3.1. LiCoO2 LiCoO2 is formed at high temperatures and has a layered structure like a-NaFeO2 based on a closepacked network of oxygen atoms with Li  and Co3 ions arranged on alternating (111) planes of the cubic rock-salt structure. It is superior in cycling behavior due to its high structural stability and can be cycled more than 500 times with 80
/90% capacity retention. In addition, its preparation is easy and capacity high. As a result, it is the most common cathode material on the market. However, cobalt is the most expensive component as compared with nickel, manganese and vanadium. In order to lower its cost and/or enhance its reversible capacity, doping is also employed, such as the doping with fluorine, magnesium, aluminum, nickel, copper or tin [89
/95]. When 1
/10 wt% of LiF are added the reversible capacity is higher than that of pure

LiNiO2 has a similar layered structure like LiCoO2. Although it is cheaper than LiCoO2 and its reversible capacity is higher than that of LiCoO2, it is difficult to prepare on a large scale with an ideal layered structure. The main reason is the difficult oxidation to Ni3. In addition, LiNiO2 decomposes easily into Li-deficient compounds Lid NiO2d (0 B/d B/1) at high heat-treatment temperatures, which act as impure phases and are unfavorable for lithium intercalation and de-intercalation. In order to overcome these disadvantages for synthesis, doping is a promising method. There are many kinds of dopant elements such as Mg, Al, Ti, Mn, Fe, Co, Zn, Ga, Nb and F. In the case of substitution by magnesium Ni2 is mainly substituted at small added amounts of Mg2, this results in good cycling behavior. With large added amounts magnesium may substitute Ni4 resulting in a quite different electrochemical performance [96,97]. Aluminum can be uniformly doped into LiNiO2 [97
/ 103]. LiAlx Ni1x O2 can be a single-phase compound with layer structure up to x /0.25, which is formed under oxygen atmosphere at 750 8C. Since Al3 is inert, disruption of the LiNiO2 structure under overcharge is prevented. Resistance for diffusion of charge carriers is lower, the diffusion coefficient for Li  increases [99]. Exothermal reactions are markedly suppressed during the charge process, the stability of the electrolyte solution increases apparently [100]. Reduction potential also shows an increase by about 0.1 V after doping with aluminum, and thus a third plateau corresponding to lithium intercalation will not appear during the normal cycling up to a final voltage below 4.3 V (prior to substitution this proceeds around 4.23 V). Only the first and second plateaus appear, whose potentials are 3.73 and 4.05 V (3.63 and 3.93 V prior to substitution) [101]. As a result, cycling behavior and tolerance to over-charge improve [97
/103]. Under oxygen flow at 700 8C LiAl0.25Ni0.75O2 can also be prepared by static electron sputtering deposition [102]. Ti4 can be doped into LiNiO2 to form layered LiNi1x Tix O2 (0.025B/x B/0.2) by solid-state reaction

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at 750 8C under an oxygen stream. It is a highly ordered and single-phase layered compound. Substituting with Ti4 ions preserves the structural integrity of the material, and impurity Ni2 ions are prevented from migrating into lithium sites [104,105]. The reversible capacity can be as high as 240 mA h g1 and cyclability is excellent (over 100 cycles) in the range 4.3
/2.8 V at C/ 5 rate [104]. Partial substitution of nickel by manganese is effective in preventing abrupt changes of the lattice constants. Crystal parameter c and Ni
/O distance increase with x , resulting in suppression of the phase transitions (hexagonal(r) 0/monoclinic(r) 0/hexagonal). The charge compensation in charge/discharge cycling of Li1x Ni0.8Mn0.2O2 could be attributed, based on results of XPS, XANES and ESR, to change valence of nickel. However, exchange of nickel into the 3a-type lithium site makes it difficult for the lithium ion to move through the lithium layer, and decreases the reversible capacity [106]. Therefore, electrochemical performance becomes poor [107]. After doping with Fe3 the potential for lithium deintercalation increases resulting in more difficult oxidation of Ni3. In addition, numerous Ni2- or Fe3ions occupy lithium sites, and thus electrochemical performance deteriorates [108]. In order to stabilize Ni4, cobalt can partly substitute nickel [105,109
/117] resulting in LiNi1x Cox O2, whose reversible capacity can be above 180 mA h g1. Ni/O and Ni/Ni distances decrease with a decrease of x in Lix NiO2 (x 5/0.8), and local distortion of the NiO6coordination sphere decreases with an increase in the doping amount of cobalt [112]. The added cobalt exists in a Co3O4 spinel structure in Li1x Ni0.85Co0.15O2 and suppresses the decomposition of Li1x NiO2 into a rocksalt phase [111]. At x /0.26 reversible capacity in the first cycle of LiNix Co1x O2 is between 205 and 210 mA h g1 depending on the particle size [113,117]. Under charging and discharging at C/2 reversible capacity can be 157 mA h g1. Its high rate capability is comparable with that of LiCoO2 [113]. Using spherical particles of Ni1x Cox (OH)2 (x/0.1, 0.2 and 0.3) as precursors, LiNi1x Cox O2 can be prepared. As compared with the Li1x Cox Ni1x O2 prepared by mixing Co(OH)2 and Ni(OH)2 its degree of disorder decreases and the ratio of the crystal parameter c /a increases, and thus electrochemical stability increases as shown in Fig. 7 [105]. Doping with cobalt can also be performed on the basis of a sol
/gel method. For example using polyvinyl alcohol as a carrier for the sol
/gel method, Lix Ni0.85Co0.15O2 can be prepared, and its crystallinity is very good when heat-treated at 600 8C and its high rate capability is better than those of materials prepared via a solid phase method [115]. Ga-doped LiNiO2 has a single hexagonal structure, other compounds such as LiGaO2 do not exist. During

3499

Fig. 7. Cycling behavior of LiNiO2 and LiNi(1xy )Mx My? O2 cathode materials for M, M?/Ti, Co, or Co/Mg in coin cells with lithium metal counter electrode measured at 60 8C between 4.1 and 3 V in 1 M LiPF6 solution of PC/EC/DMC(1/1/3) (from reference [105]).

the charging process, the hexagonal structure is maintained, a monoclinic phase or other types of hexagonal structures are not observed. The crystal parameters change slowly and continuously. Within the range of 3.0
/4.3 V reversible capacity is above 190 mA h g1, and capacity retention is above 95% after 100 cycles. When the charge voltage is higher, 4.4 or 4.5 V, reversible capacity is above 200 mA h g1 and capacity retention is also very good. It is stable during overcharging [97,118]. Doping with arsenium, calcium, indium or niobium shows only very limited improvements in cell performance [119]. After partial substitution of oxygen with fluorine Xray diffraction pattern reveals that the crystal structure still changes during charge and discharge. However, its cycling behavior improves, which is mainly due to a decrease of the inner electrical resistance [120]. In addition, other reports indicate that this kind of substitution also suppresses the phase transition and thus cycling behavior improves [121]. Of course, two or more different types of heteroatoms can be doped into LiNiO2 [105,121
/128]. The reversible capacity of Li1x Ni1x Coy O2z Fz doped with both cobalt and fluorine via solid-state reaction can be 182 mA h g1. After the first 100 cycles capacity fades only 2.8%. In following cycles, capacity fades much less. This can be explained by invoking the improvements caused by substitution with cobalt and fluorine [123]. The reversible capacity of LiNi0.75Ti0.125Mg0.125O2 and LiNi0.70Ti0.15Mg0.15O2 can be up to 190 mA h g1. In the case of LiNiO2 during thermal studies an exothermal peak is found already at 220 8C. When LiNi1x Tix /2Mgx /2O2 is in its charged state the extent of the exothermal process at 220 8C declines with an increase in x [124]. No exothermal peak is observed in thermogravimmetric experiments below 400 8C with the compounds in the charged state when x ]/0.25 [124].

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The optimum composition for co-doping with cobalt and aluminum is Li(Ni0.84Co0.16)0.97Al0.03O2. The reversible capacity can be 185 mA h g1, irreversible capacity in the first cycle is only 25 mA h g1, and cycling behavior is very good. In addition, thermal stability improves strikingly [125]. The stability of LiNiO2 doped with strontium and cobalt, viz. LiSr0.02Ni0.9Co0.1O2, is insufficient, but a coated MgO layer on its surface can improve stability. After the surface coating, the initial thermal decomposition temperature is increased by about 10 8C, and the exothermal process is sharply suppressed. High rate capability such as at 1C markedly improved, the possibility for explosion declines [126]. Doping with cobalt and manganese can yield Li8(MnCo2Ni5)O16, which is a single-phase layered structure. Its reversible capacity is 150 mA h g 1, and capacity fading is only 0.41 mA h g1 per cycle [127]. LiNiO2 doped with zinc and aluminum on the basis of a sol
/gel method can result in a reversible capacity as high as 245 mA h g1; there is no capacity fading in the first ten cycles [128]. These results show again that there are two kinds of effects from the introduced heteroatoms: favorable and unwelcome. Unfavorable effects are from the occupation of Li  sites by the dopants such as Mn3, Fe3, Ca2, In3, Nb3 [106,108,119], and thus lithium intercalation and de-intercalation are hindered, electrochemical performance deteriorates and voltage increases. Favorable factors include the following: (1) substituting Ni2 in the impure phase and inhibiting side action from Ni2 [96,97], (2) forming inert layered LiMO2 (M /Al, Ga and Co) and substituting part of LiNiO2 to prevent destroying of LiNiO2 structure from over-charge [97
/102,118], (3) increasing the stability of the crystal structure to circumvent phase transformation [104,111,118,121], (4) decreasing charge transfer resistance and increasing Li  diffusion coefficient [120], (5) blocking Ni2 from diffusing into Li  site [104]. Consequently, reversible capacity enhances, up to 245 mA h g1, exothermal amount decreases, high rate capability and over-charge behavior improve. Of course, co-actions of two or more kinds of favorable dopants will result in better electrochemical performance. From this point of view, doped LiNiO2 can be a good candidate as a cathode material for lithium ion batteries.

3.3. LiMn2O4 LiMn2O4 has a spinel structure, different from LiCoO2 and LiNiO2. Since it is cheap and has small environmental impact, it is fairly attractive. However, its capacity fades slowly, and this prevents its commercial use. This fading is mainly due to the following three factors:

1) Dissolution of Mn3: At the end of discharge, the concentration of Mn3 arrives at its highest level. The Mn3 at the surface may disproportionate as the following equation 9 [1,129
/133] 4 2 2Mn3 solid 0 Mnsolid Mnsolution

(9)

Mn2-ions from this disproportionation dissolve in electrolyte solutions. 2) Jahn
/Teller effect: at the end of discharge, the Jahn
/Teller effect happening at first on the surface of several particles may expand into the overall composition of Li1d [Mn2]O4. Thermodynamically speaking this system is not really at equilibrium [130,132]. The phase transition from cubic into tetragonal symmetry is a first-order process. Even though this kind of change is small, it is big enough to destroy the structure to form a tetrahedral structure, which is low in symmetry and high in disorder [134]. 3) In organic solvents, the highly de-lithiated particles are not stable at the end of discharge, viz. the high oxidation ability of Mn4 will lead to decomposition of the solvents [132,135,136]. These factors may be active at the same time resulting in a fading capacity. If the spinel structure can be modified, at least part of these phenomena may be suppressed. One of the effective methods is doping with heteroatoms. A lot of elements have been introduced into LiMn2O4 to investigate doping effects [129]. Following doping with cations and anions is reviewed. 3.3.1. Doping with cations Cations of Li, B, Mg, Al, Ti, Cr, Fe, Co, Ni, Cu, Zn, Ga and Y have been investigated as possible dopants. Introduction of lithium can be performed via two methods. An excess of lithium salt is added during the synthesis of the spinel Li[Mn2]O4 resulting in Li1x [Mn2]O4 (x /0) [137]. The other method involves reaction of Li[Mn2]O4 with lithium butyl to form Li1x [Mn2]O4 according to [138]: Li[Mn2 ]O4 xLiC4 H9 0 Li1x [Mn2 ]O4 1=2C8 H18

(10)

In the former case, the structure of Li1x [Mn2]O4 changes with heat-treatment temperature T and value of x . When x B/0.14 and T is 700 8C, a single spinel structure is obtained. When T /750 8C, the tetrahedral spinel changes into a rhombic one and decomposes forming LiMn2O4 and Mn3O4. 3LiMn2 O4 0 3LiMn2 O4 Mn3 O4

(11)

The produced LiMn2O4 is not stable and disproportionates according to: 3LiMn2 O4 1=2O2 0 LiMn2 O4 Li2 MnO3

(12)

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and forms rock-salt structure Li2MnO3. At low temperatures and with x /0.14, Li2MnO3 will also form. The initial reversible capacity of Li1x [Mn2]O4 synthesized at 750 8C is lower than that of LiMn2O4, but the cycling behavior is better. The average reversible capacity in the first 50 cycles is over 120 mA h g1. Li1x [Mn2]O4 synthesized in the latter way is a mixture of LiMn2O4 and Mn3O4 [138]. When charged to the plateau of about 3 V, the lithium introduced via chemical reaction can be utilized completely, it can actually reimburse some of the capacity loss caused by the anode in the first cycle after manufacturing of a lithium ion battery, and enhance the practical capacity. Furthermore, fading of capacity becomes slower. The ionic radius of trivalent boron is 27 pm, much smaller than that of trivalent Mn3 (65 pm). When it is introduced into LiMn2O4, it prefers trigonal or tetrahedral coordination sites, and the spinel structure is disrupted. After addition of B2O3 the vacancies among particles and grains’ porosity decrease sharply. As a result, electrochemical properties deteriorate, and the initial reversible capacity is low ( B/50 mA h g1) and capacity fades very quickly [139,140]. The effect of adding magnesium into LiMn2O4 is the same as that of an excess of lithium, viz. enhancing the average oxidation state of manganese and impeding Jahn
/Teller distortion. Using lithium metal as a reference, there is no capacity fading after 20 cycles and capacity is above 100 mA h g1. The reversible capacity cycled between 4.3 and 1.6 V can be up to 180 mA h g1, but it declines with cycling [141]. The radius of trivalent Al3 is 53.5 pm, smaller than that of trivalent Mn3 [140,141]. When it is introduced into spinel LiMn2O4, Al3 is situated at tetrahedral 3 sites, the lattice contracts forming [Al3 2 ]T[LiAl3 ]OO8 structure (T and O indicating tetrahedral and octahedral sites). Consequently, in the LiAl0.02Mn1.94O4 spinel structure Al3 may substitute for Li  in the tetrahedral sites 8a, and may cause the original Li  to move to octahedral sites 16c. Li  at the octahedral sites cannot de-intercalate at 4 V. As a result, the disordering degree of cations increases and electrochemical performance declines. This is different from the doping of aluminum into LiCoO2 [94,95] and LiNiO2 [97,111]. Other results indicate that the reversible capacity of LiAlx Mn2O4 decreases slightly when x 5/0.05, and its cycling behavior improves much. There is no capacity fading in the first 30 cycles [129,139,142,143] since the introduction of Al3 results in high average oxidation state of manganese ions. In addition, the cycling behavior of LiAlx Mn2x O4 in the 3 V region can improve if suitable precursors are selected [144]. The radius of Ti3 is 67 pm, it is easily incorporated into the lattice structure of LiMn2O4. However, it can be oxidized easily into Ti4, and this results in an average valence of manganese below 3.5. During charge and

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discharge the Jahn
/Teller structural distortion happens to cause rapid degradation in physicochemical properties and a resulting capacity decline [140]. Its action is quite different from that in LiNiO2 [104]. The radius of Cr3 is 61.5 pm, very near to that of Mn3. It exists in a stable d3 configuration and prefers an octahedral coordination [140,145
/147]. Therefore, in the composite oxide LiCrx Mn2x O4, it is still a singlephase of spinel even if x is as high as 1/3. Lattice parameters decrease linearly with x in LiCrx Mn2x O4 [145]. During the charging process, the cubic symmetry of this spinel is not destroyed, and its cycling behavior improves much. Of course, with an increase in the doping amount of chromium, capacity will decrease, and sometime sharply, which can be partly attributed to lattice contraction [145]. The optimum composition is one in which 0.6% Mn3 is substituted by Cr3. Its initial capacity decreases just by 5
/10 mA h g1, and cycling behavior is much improved (Fig. 8). After 100 cycles, reversible capacity can still be 110 mA h g1. This improvement in cycling behavior is mainly due to an increase in the stability of the spinel structure, which can also be indirectly derived from the binding energy in the corresponding metal oxides MO2. The binding energy of M
/O in MnO2 (a-type) and CrO2 is 946 and 1029 kJ mol 1. Concomitantly, when a spinel structure is good in stability, the dissolution according to (Eq. (9)) will be suppressed [129]. Although the radius of Fe3 is 64.5 pm, near to that of Mn3, it prefers a high-spin d5 configuration. Consequently, it exists in an inverse spinel structure LiFe5O8 and this can easily result in cation disorder. Therefore, the coulombic efficiency is not high and capacity fades fast. In addition, Fe3 may catalyze the decomposition of electrolyte solution [140]. If the oxidation of Fe3 to Fe4 is utilized, LiFex Mn2x O4 spinel can be used as 5 V cathode material for lithium ion battery [148].

Fig. 8. Cycling behavior of cathode materials of LiMn2O4 and doped Lix M1/6Mn11/6O4: M/Cr, Co and Ni (from reference [141]).

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Cobalt exists in the spinel structure LiCoy Mn2y O4 as trivalent cation [146,149
/151]. Similar to the doping with Cr3, its doping increases the stability of the spinel structure (binding energy of Co
/O in CoO2 is 1142 kJ mol 1). During the charge and discharge process, the change in volume is minor ( 5/5%), and thus the spinel structure is not easily disrupted. In addition, the conductivity of LiCoy Mn2y O4 is much higher than that of LiMn2O4 because the diffusion coefficient of Li  increases from 9.2 /1014
/2.6 /1012 to 2.4 / 1012
/1.4 /1011 m2 s 1 (measured at various states of charge). Furthermore, after doping with cobalt, the particle size becomes larger and the specific surface area decreases, consequently, the contact area between active materials and electrolytes also becomes smaller, and the decomposition rate of electrolytes and self-discharge rate decrease. The presence of cobalt inhibits the passivation process occurring on the cathode surface, increases the exchange current density and facilitates the charge-transfer reaction of the active material [150]. All of these factors can favor the reversible intercalation and de-intercalation of lithium, and result in evident improvement in cycling behavior (Fig. 8). From its capacity and cycling behavior, it can be seen that this spinel can be used not only as cathode material of 4 Vlithium ion battery (4.2
/3.7 V) [146,150] but also as that of 3 V-lithium ion battery (3.3
/2.3 V) [149]. By the way, if the charge voltage is increased to 5.3V, a new plateau centered at about 5.0 V is observed and the Co-doped LiCo(1x )/2Mn(3x )/2O4 (/15/x 5/1) can be a cathode material for 5 V lithium ion battery [152]. Nickel exists as divalent cation in LiMn2O4 [146]. Although the intercalation of lithium decreases the average valence of manganese to values below 3.5, sometimes as low as 3.3, distortion of the tetrahedral phase is not observed [153]. Its action is similar to that of cobalt and chromium, which can stabilize the octahedral spinel sites (binding energy of Ni
/O in NiO2 is 1029 kJ mol1), and it improves its cycling behavior (Fig. 8). When the charge voltage increases from 4.3 to 4.9 V, there is a new voltage plateau around 4.7 V, which corresponds to the change from Ni3 to Ni4, and can be used as cathode material for a 5 Vlithium ion battery [154]. The introduction of Ni2 into spinel structure can yield Li[Mn1.5Ni0.5]O4 [153], where lithium can intercalate. The intercalation at 3 V is a twophase reaction, the end product of lithiation is a stoichiometric rock-salt structure Li2[Mn1.5Ni0.5]O4. The synthesis temperature for Ni-doped spinel lithium manganese oxides should not be too high. When T / 650 8C, the phase Lix Nix O will also appear resulting in deterioration of electrochemical performance. The capacity of LiNi0.5Mn1.5O4 prepared at 600 8C under cycling between 4.9
/3.0 V can be stable above 100 mA h g1 [154].

When copper is introduced into spinel LiMn2O4, it exists in di- and tri-valent forms [155,156], the chemical III III,IV formulae can be written as LiCuII x Cuy Mn2xy O4. At around 4.9 V there is a new voltage plateau as show in Fig. 9, corresponding to the oxidation/reduction of the Cu2/Cu3 couple. This material can be used as cathode material for 5 V-lithium ion batteries. Like the doping with other elements such as Cr, Co and Ni, reversible capacity decreases but cycling behavior improves. The inclusion of copper in the spinel structure enhances the electrochemical stability during cycling. The potential for the oxidation of Mn3 to Mn4 shifts slightly to a higher value. The compound exists in a single-phase although there is an impure lithium
/copper oxide phase. The substitution of manganese by copper enhances the reactivity of the spinel structure towards hydrogen. After the introduction of zinc into the spinel structure, there is no Jahn
/Teller distortion since Zn2 is in 3d10 configuration. Similar to the effects of doping with lithium and magnesium, the Jahn
/Teller effect is blocked and cycling behavior improves. The reversible capacity of LiZn0.05Mn1.95O4 can stay at 102 mA h g1 after 20 cycles [132]. Regarding doping with gallium, there are two different reports. The radius of Ga3 (62 pm) is close to that of Mn3. Like Al3 it can form an inverse spinel structure, LiGa5O8, and thus leads to disordering of the lattice structure and fast fading in capacity [140]. In other studies it was concluded that Ga-doped LiGax Mn2x O4 exists in a single spinel phase, and the cubic symmetry is maintained [139,157,158]. As a reason of the similarity of Ga3 to that of Zn2 its configuration, viz. 3d10, is invoked, there is no Jahn
/Teller distortion and the crystal parameter a (82.27 pm) is almost the same. The ratio of Mn3/Mn4 is below 1 and the distortion from Jahn
/Teller effect during charge and discharge process decreases [157]. Ga3 ions selectively occupy the octahedral (16d) sites of the lattice. There are no gallium ions situated at the

Fig. 9. Voltage-capacity curves obtained during the third cycle for LiCux Mn2x O4 (05/x 5/0.5) in steps of x /0.1 (from reference [156]).

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tetrahedral (8a) or the additional octahedral (16c) sites [158]. As a result, the cycling behavior is improved. Like doping with other elements, there is a decrease in capacity. When x/0.05 in LiGax Mn2x O4, its performance is optimal, and there is almost no decrease in capacity. Furthermore, its cycling behavior is good. After the introduction of trivalent yttrium into LiMn2O4, the obtained material exists in a two-phase structure. Consequently, its electrochemical performance is not satisfactory [140]. The above results of cations doping show that there are two kinds of effects: favorable and unwelcome. Favorable doping is due to the following actions: (1) increase in Mn valence suppressing Jahn
/Teller effect [137,138,141,142], (2) improvement in stability of spinel frame of Li[Mn2]O4 lowering structural change during charge and discharge process and inhibiting the solution reaction of Mn happening as equation of Eq. (9) [140,145
/147,149
/153], (3) increase in conductivity favoring reversible lithium intercalation and de-intercalation [150], (4) decrease of surface area resulting in decrease in contacting area between active cathode materials with electrolyte and decomposition rate of electrolyte and self-discharge rate [150]. Unfavorable effects of doping are due to the following factors: (1) decrease in Mn valence promoting Jahn
/Teller distortion [140], (2) formation of heterogeneous phase as impurity destroying the uniformity of spinel frame structure and unfavorable for lithium intercalation and de-intercalation [140], (3) decrease in volume of crysallite unit of spinel Li[Mn2]O4 inhibiting movement of lithium [139
/141]. Of course, as to favorable doping, it does not mean more content of dopant will achieve better electrochemical performance. When it is too much, it will also result in the production of heterogeneous phase and deterioration of electrochemical performance. By the way, there is usually a trade-off between cycle life and capacity in the case of these favorable dopants. In addition, doping of LiMn2O4 produces new materials such as 5 V cathode materials. 3.3.2. Doping with anions Anions of oxygen, fluorine, iodine and sulfur can be used as dopants. By changing the synthesis conditions such as precursor materials, heat-treatment temperature and soak time, the amount of oxygen in LiMn2O4 can be changed [159
/162]. Mixing MnCO3 and Li2CO3 in molar ratio of 2:1 and reacting in a temperature range 300
/400 8C by a solid-state reaction LiMn2O4d is obtained, and the maximal component can be Li2Mn4O9. LiMn2O4d is a defect spinel, the cations arrange in (Li0.89I0.11)[Mn1.78I0.22]O4 (represents vacant positions) and the valence of all manganese ions is /4. When lithium initially intercalates, there is only one voltage plateau, which is at 3 V. But in the following cycles, intercalation

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and de-intercalation of lithium happen at 3 and 4 V plateaus, correspondingly lithium intercalates into and de-intercalates from octahedral and tetrahedral sites [162]. Jahn
/Teller effect will happen at the end of discharge, viz. Li4Mn4O9 (c /a/1.14), where the average state of oxidation of manganese ions is 3.5. Fluorine can substitute partially oxygen and forms Li1x Mn2x O4y Fy (0 B/y B/0.5). Since the electronegativity of fluorine is bigger than that of oxygen, its electron-attracting ability is larger, this decreases the solubility of manganese ions in organic solvents. The storage stability at relatively high temperature (about 50 8C) increases [163
/165]. However, after partial substitution of oxygen by fluorine, the average oxidation state of manganese is lowered. In order to compensate this effect, the amount of lithium in this component should increase, viz. x should be a little bigger so as to ensure an average oxidation state of manganese above 3.5. On this basis, doping with aluminum can be further performed to improve the stability at high temperature [165,166]. Reduction of anhydrous NaMnO4 with LiI in anhydrous acetonitrile can get I-doped LiMn2O4 [167
/169]. In fact, the prepared material is evidently not of a spinel structure, and should be amorphous. Its ratio of I/Na is dependent on that of the reactants. Since its conductivity is lower than that of spinel LiMn2O4, initial electrochemical performance is not good, and capacity fades. When it is mixed with conductive carbonaceous materials by ball-milling, electrochemical performance improves considerably. When charged at 0.05, 0.5 and 1 mA cm2, the maximum capacity can be 335, 275 and 220 mA h g1. The amorphous nature results in less dramatic change in structure upon insertion of Li , and the stress due to lithium intercalation is small. As a result, its cycling stability increases (Fig. 10) and it is a potential material with high rate capability [168]. The product of partial substitution of oxygen by sulfur, LiMn2O3.98S0.02, can be prepared by a sol
/gel method. Its initial capacity is only 80 mA h g1, and increases in the following cycles, and reaches 99 mA h g1 after 20 cycles. Since sulfur atoms are bigger than oxygen ones, the structure stability can be kept during cycling and it can circumvent the Jahn
/Teller distortion around 3 V [170]. On this basis, cations such as Al3 and Mg2 can be further doped [171
/174]. For example, the cubic spinel structure of Li1.03Al0.2Mn1.8O3.96S0.04 does not change during cycling [171]. LiAl0.24Mn1.76O3.98S0.02 will not show Jahn
/Teller effects at both voltage plateaus of 3 and 4V (2.4
/4.3 V), either, and reversible capacity can be up to 215 mA h g1 [172]. The excellent cycling behavior of the oxysulfide spinel at high temperature over 4 V region is due to small structural degradation of this material [174]. In short, favorable doping of anions is concentrated on the following two aspects: (1) stabilization of the

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cycles, reversible capacity can be above 410 mA h g1. Recent studies show that the lithiation of Cu0.1V2O5 xerogel cathode proceeds with the concurrent formation of copper metal, and the starting parent compound can be regenerated upon removal of lithium, which is attributed to the excellent structural and electronic reversibility [180]. Results concerning doping with heteroatoms resulting in the formation of inverse-spinel structures such as LiNiVO4 and the effect of multiple anions such as PO3 4 have been reviewed elsewhere [181,182]. As a conclusion, favorable doping for V2O5 should mainly focus on the improvement of mechanical structural stability. Otherwise its poor cycling behavior will make it useless for commercial consideration.

4. Summary Fig. 10. Cycling behavior of amorphous Li1.51Na0.51MnO2.85I0.12 ballmilled with 25 wt.% fine carbon for various time: (a) 5, (b) 10, (c) 20, and (d) 40 min. The data were recorded with a current density is 0.5 mA cm 2 (from reference [168]).

spinel structure by e.g. fluorine, (2) reduction of volumetric change or stress from lithium intercalation such as I and S. This is useful for consideration of doping with other anions. 3.4. V2O5 Vanadium is also cheap and easily derived from existing mineral deposits like Mn, and its oxides are attractive cathode materials due to accommodation of three stable oxidation states (V5, V4 and V3) within its closely packed oxygen structure. So far, a wide range of vanadium oxides has been investigated. a-V2O5 has a layered structure with a distorted close-packed oxygen array. Its theoretical capacity is the highest within the vanadium oxide family (442 mA h g1) and three lithium ions can be inserted reaching a composition of Li3V2O5. However, V2O5 prepared by normal sol
/gel methods is still not satisfactory since its capacity fades fast. To modify its properties, one possibility is to introduce heteroatoms such as iron, aluminum, chromium, copper and cesium during the sol
/gel preparation, like doping of LiMn2O4 [175
/179]. The dopant elements are distributed homogeneously in the prepared products and they increase interaction between V2O5 layers. As a result, the stability of the layered structure increases during charge and discharge showing better cycling properties. For example, in the case of Cr0.11V2O5.16 (V2O5 derived structure), the formation of short CrO6 octahedral chains link V2O5 layers to increase its three-dimensional stability [179]. After 40

The introduction of heteroatoms greatly influences the structure and the electronic state of electrode materials. There are two kinds of effects: favorable and unwelcome. The specific effects depend on the species and states of the heteroatoms and the parent electrode materials. When applied in lithium ion batteries, advantageous factors should be exploited and the unfavorable side effects should be avoided as much as possible. For example, when the chemical states of the doped elements are optimized, electrochemical performance of carbonaceous materials is much improved. With tin-based oxide anode materials, the introduced heteroatoms will favor the formation of amorphous glass structure and lithium diffusion. The introduction of inert heteroatom of a ductile material will reduce markedly the volumetric change of alloy anodes during cycling. Introduction of heteroatoms into LiCoO2 will not only lower the cost but also improve electrochemical performance. The introduced heteroatoms can stabilize Ni4 in LiNiO2, remove Ni2 in impure phase, enhance reversible capacity and improve cycling behavior. As to spinel LiMn2O4, the favorable heteroatoms can make the spinel structure more stable, suppress the happening of Jahn
/Teller distortion, and block the decomposition of electrolyte solution during cycling. In the case of V2O5, the heteroatoms will increase the mechanical stability of V2O5 layers. So far, great advances towards improvement of electrochemical performance of the electrode materials have been achieved. Of course, still some problems need investigation especially from a theoretical point of view [183,184]. With the understanding of the molecular structure of electrode bodies, these kinds of phenomena can be more clearly explained [185]. Simultaneously, this will propel the advance of composite/doping technology to prepare better and novel electrode materials.

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Acknowledgements Financial support from Alexander von Humboldt Foundation (Y.P. Wu) is appreciated. Part of this work is based on previous research financially supported by China Natural Sciences Foundation and China Postdoctor Foundation. A. Manthiram, P. Novak, M. M. Thackeray and M. Winter provided reprints and valuable further information.

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