Nano Today (2011) 6, 28—41
available at www.sciencedirect.com
journal homepage: www.elsevier.com/locate/nanotoday
Roles of nanosize in lithium reactive nanomaterials for lithium ion batteries Kyu Tae Lee, Jaephil Cho ∗ Interdisciplinary School of Green Energy and Converging Research Center for Innovative Battery Technologies, Ulsan National Institute of Science and Technology (UNIST), Ulsan 689-798, Republic of Korea Received 6 October 2010; received in revised form 3 November 2010; accepted 4 November 2010 Available online 24 December 2010
KEYWORDS Nanosize; Lithium ion batteries; Electrode; Nanostructured materials; Electrochemistry
Summary There have been exciting developments in new electrode materials for lithium ion batteries in the past decade. Nanostructured materials have emerged as highly suitable candidates, including Si anode and LiFePO4 cathode materials, and are starting to be used in the marketplace. Promising electrochemical properties including excellent kinetics and cycling stability have driven research interest in nanomaterials. This review highlights the major roles of nanosize in lithium reactive nanomaterials for Li ion batteries, with the aim of providing nanomaterial scientists with a better understanding of electrochemical concepts that can be exploited for tailored design of electrode materials. © 2010 Elsevier Ltd. All rights reserved.
Introduction Lithium ion batteries have been the focal point of much research due to their higher energy density compared to existing electrochemical systems such as Ni-MH batteries. On the basis of this strength, Li ion batteries have been successfully used in various applications such as hybrid electric vehicles, mobile electric applications, and renewable energy storage devices. Lithium ion batteries are comprised of three major components: a cathode, anode, and electrolyte (Fig. 1) [1,2]. During electrochemical reactions, Li ions move from the cathode to the anode through the electrolyte or vice versa, and thus there are three reaction
∗ Corresponding author. Tel.: +82 52 217 2910; fax: +82 52 217 2909. E-mail address: [email protected]
steps: Li ion diffusion within solid state electrode materials, a charge transfer reaction at the interface between the electrode and electrolyte, and Li ion movement in the electrolyte. Among these, the solid state diffusion of Li ions is generally considered to be the rate-determining step of Li ion batteries. Therefore, nanostructured materials have been extensively explored in efforts to enhance kinetic properties by decreasing the diffusion length to a nanometer scale . However, the solid state diffusion of Li ions is not the sole major factor affecting the kinetics of electrochemical reactions, and thus the electrochemical system of Li ion batteries is more complicated. Recently, two notable reports have attempted to explain the kinetics of cathode materials in Li ion batteries. One study explained the kinetics using a domino-cascade model of the reaction mechanism for LiFePO4 , and the other attributed the kinetics to enhanced rate performance of LiFePO4 by surface coating with ionic conductive layers. Li de/insertion in most LiFePO4 materials excluding highly
1748-0132/$ — see front matter © 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.nantod.2010.11.002
Nanosize in lithium reactive nanomaterials
Figure 1 Schematic representation of a lithium-ion battery. Various materials with different structures can be used for anode and cathode materials.
defective LiFePO4  proceeds through a two-phase reaction, because the nucleation rate is faster than solid state diffusion in a particle . It means two phases can be observed in a single particle, as veriﬁed by TEM and EELS analyses [5,6]. However, Delmas et al. proposed that the nucleation rate is the rate-determining step in the case of LiFePO4 nanoparticles, and termed this the dominocascade model . High-resolution TEM and XRD studies on partially delithiated nanoparticles (<100 nm in diameter) have revealed that the individual particles are single phase (either Li-rich Li1−y FePO4 or Li-poor Lix FePO4 ), contrasting with the bulkier samples reported previously wherein two phases of Li1−y FePO4 and Lix FePO4 were clearly exhibited in a single particle. This phenomenon is attributed to the solid state diffusion step of Li ions completing as soon as the nucleation reaction happens. This explanation is not clear, however, because it does not consider various factors related to the thermodynamics of electrochemical systems, such as heterogeneous particle size distribution . Nevertheless, this proposal provides a good opportunity to reconsider the kinetics of Li ion batteries at nanoscale dimensions. If the domino-cascade model is correct, it would be pointless to try to make smaller LiFePO4 materials, and dimensions of tens of nm would be small enough to exhibit charge transfer limiting reactions. Therefore, the dominocascade model suggests that the strategy for manipulating LiFePO4 nanoparticles should entail surface modiﬁcation to decrease the charge transfer resistance between the electrode and electrolyte. A similar proposal was also made by the Ceder group, stirring considerable debate [9—11]. They showed good rate performance of LiFePO4 nanoparticles by surface modiﬁcation, and they contended that the charge transfer resistance is one of the dominant resistances. This coincides well with the domino-cascade model, as both consider the charge transfer reaction to be the rate-determining step. The aim of the present review is not to determine the validity of the above explanation. Rather, the focus of this
work is to demonstrate that nanosize electrode materials can induce various effects on the reaction mechanism as well as the performance of Li ion batteries, although the original aim of pursuing nanoscale dimensions was merely to provide a short diffusion path. Most previous studies on nanomaterials for Li ion batteries focused on decreasing the diffusion length by reducing particle size . However, this is just one of a number of nanosize effects. Nanomaterials can alter electrochemical conditions, and different strategies should thus be adopted to enhance electrochemical performances relative to values provided by bulkier materials. Therefore, it is necessary to attain a comprehensive understanding of the effects of nanosize in order to establish new strategies. This review thus presents the various roles of nanosize in lithium reactive nanomaterials for lithium ion batteries.
Discussion Diffusion length effect As noted above, nanomaterials are beneﬁcial for Li ion batteries as a result of providing short diffusion length [13,14]. Notably, LiMPO4 (M = Fe, Mn) olivines are promising cathode materials for EV applications due to enhanced safety; however, they have poor ionic conductivities, resulting in poor electrochemical performances of bulk materials. In response, numerous nanosized LiMPO4 (M = Fe, Mn) materials have been synthesized through various synthetic routes. LiMPO4 materials have a one-dimensional ionic diffusion path along the [0 1 0] direction based on the space group of Pnma [15—19]. This means the length of particles along the b direction should be decreased so as to obtain a short diffusion path. A smaller dimension along the [0 1 0] direction can be obtained via the realization of three types of morphologies: (1) zero-dimensional spherical nanoparticles; (2) one-dimensional nanorods oriented with long axes along the a or c direction; or (3) two-dimensional nanoplates with
K.T. Lee, J. Cho
Figure 2 (a) SEM image, (b) TEM image, (c) charge—discharge curves and (d) the rate behavior of 2D LiMnPO4 . Reprinted with permission from .
a (1 0 1) basal plane. 0D spherical nanoparticles and 1D nanorods of LiFePO4 olivine materials can be synthesized by conventional solid state synthetic methods [20—22] and polyol synthesis [8,23,24], respectively. Also, by employing hydrothermal synthesis, 2D nanoplate morphologies can be obtained [25,26]. It should be noted that the above relationship between morphologies and synthetic methods is not strict, and there are many exceptional cases. For example, LiMnPO4 nanoplates have been synthesized by both a polyol method  (Fig. 2) and a solid state reaction in molten hydrocarbon with self-assembly . These nanostructured LiMPO4 materials exhibited better rate performances than bulkier materials due to smaller diffusion length (Fig. 2). Another alternative to commercialized LiCoO2 cathode materials is LiMn2 O4 spinel materials (space group: Fd-3m). These materials have also been considered for use in EV applications owing to their good safety [29,30]. The electrode materials for EVs should have not only a desired level of safety but also good rate performance to charge or discharge batteries quickly. For this reason, nanostructured LiMn2 O4 spinel materials have been synthesized. Representative nanostructures include ordered mesoporous materials with a wall thickness of a few nm. Ordered mesoporous LiMn2 O4 spinel materials were obtained by hard template synthesis using ordered mesoporous silica (KIT-6), as shown in Fig. 3a [31,32]. Ordered mesoporous manganese oxides were prepared by sequential steps of (1) inﬁltration of pre-
cursors in KIT-6; (2) heating precursor-inﬁltrated KIT-6; and (3) dissolution of KIT-6 in NaOH solution. The Li precursor is then inﬁltrated in mesoporous manganese oxides, or Li ions are inserted in mesoporous manganese oxides by reduction with LiI. Finally, ordered mesoporous LiMn2 O4 materials are crystallized into a spinel structure by heat treatment at 350 ◦ C. Their wall thickness was about 7 nm, resulting in a very short diffusion path. They also showed substantially better rate performance compared to bulkier LiMn2 O4 spinels, even at a 5C rate, and they maintained 80% of the reversible capacity at a 0.1C rate (Fig. 3). The same hard template method was also successfully adopted in the synthesis of ordered mesoporous low temperature LiCoO2 materials with a spinel structure . However, conventional LiCoO2 was synthesized at high temperature (ca. 800 ◦ C), and it crystallized into a layered ␣-NaFeO2 structure. Also, with this procedure it is still not possible to synthesize ordered mesoporous LiFePO4 olivine materials, although partially ordered mesoporous Li3 Fe2 (PO4 )3 have been synthesized by a soft template method using a cationic surfactant . Only nanowires or hollow LiFePO4 olivine materials have been obtained by the hard template method, because it is difﬁcult to inhibit growth of crystallites during the heating process at high temperature (ca. >500 ◦ C). Other examples of nanostructures are nanotubes and hollow structures. Nanotubes can also be obtained by a hard
Nanosize in lithium reactive nanomaterials
Figure 3 (a) TEM images of mesoporous (a) Mn2 O3 , (b) Mn3 O4 , and (c) Li1.12 Mn1.88 O4 . (d) High-resolution TEM image of mesoporous Li1.12 Mn1.88 O4 . (e and f) TEM images of mesoporous Li1.12 Mn1.88 O4 (e) after 20 cycles between 3 and 4.3 V and (f) after 30 cycles between 2 and 4.5 V. (b) Rate capability for mesoporous Li1.12 Mn1.88 O4 , bulk Li1.05 Mn1.95 O4 , and bulk Li1.12 Mn1.88 O4 ; capacity retention expressed as percentage capacity at 30 mA/g (0.30C). (c) Cycling data for mesoporous Li1.12 Mn1.88 O4 , bulk Li1.05 Mn1.95 O4 , and bulk Li1.12 Mn1.88 O4 between 3 and 4.3 V at a rate of 3000 mA/g (30C). The inset shows the load curves, 5th cycle, 25th cycle, and 50th cycle for mesoporous Li1.12 Mn1.88 O4 . (d) Cycling data for mesoporous Li1.12 Mn1.88 O4 , bulk Li1.12 Mn1.88 O4 , bulk Li1.05 Mn1.95 O4 , nanoparticulate Li1.12 Mn1.88 O4 , and nanoparticulate Li1.05 Mn1.95 O4 at 50 ◦ C at 30 mA/g (0.30C) between 3 and 4.3 V. Reprinted with permission from [31,32].
template method, and anodized aluminium oxide (AAO) is typically used as a template. The preparation procedure is similar to that employed for the synthesis of ordered mesoporous materials. Precursor solutions are inﬁltrated in the pores of AAO templates and then annealed at high temperature. The same steps are usually repeated several times, and AAO is then dissolved in a NaOH solution. LiMn2 O4 spinel nanotubes  and coaxial MnO2 /carbon nanotubes  were obtained with the AAO template method. In the case of hollow structures, a number of synthetic methods have been developed. Hollow LiFePO4 spheres were synthesized using a sequential precipitation reaction . Hollow LiVOPO4 spheres were also obtained by one pot hydrothermal synthesis . Both methods employ a similar strategy of using reactants as a self-sacriﬁcial template, which is different from conventional hard template methods. Hollow spheres are usually obtained by using carbon or silica sphere templates. Precursors are precipitated on the surface of the template spheres, and annealed at high temperature. The template spheres are then eliminated by dissolution of silica or burning of carbon spheres. How-
ever, in the case of the self-sacriﬁcial template for hollow LiFePO4 , Li3 PO4 is ﬁrst precipitated, and Fe3 (PO4 )2 is precipitated on the surface of Li3 PO4 by adding FeSO4 . The remaining unreacted Fe2+ in solution then reacts with PO4 3− dissolved from Li3 PO4 core precipitates due to the different solubility of Li3 PO4 and Fe3 (PO4 )2 . The precipitated hollow Fe3 (PO4 )2 is crystallized into LiFePO4 by additional hydrothermal treatment and further heating under an inert atmosphere at 700 ◦ C. Thus, the Li3 PO4 core is sacriﬁced to form a hollow structure. The formation mechanism of hollow LiVOPO4 is similar to that of hollow LiFePO4 . LiVOPO4 is formed on the surface of V2 O5 by reacting with LiOH in hydrothermal conditions at 250 ◦ C, and LiVOPO4 crystallites gradually grow on the outer surface of LiVOPO4 , accompanying the consumption of V2 O5 , resulting in a hollow structure. 1D nanorods of LiMn2 O4 spinels were also synthesized by a two step synthetic method: [39,40] (1) hydrothermal synthesis of manganese oxide nanorods, and (2) further reaction with Li precursors. 1D nanorods of manganese oxide are prepared ﬁrst by hydrothermal synthesis. LiOH is then coated on manganese oxides by mechanical slurry mixing, or Li ions
K.T. Lee, J. Cho
Figure 4 (a) (A) Photograph of a 3DOM carbon electrode attached to a nickel mesh. (B) Cross-sectional representation of an interpenetrating electrode and electrolyte layers in the cell. (b) SEM images of (A) 3DOM carbon before and (B) after potentiostatic EP with PPO. (C) SEM image of the sample shown in (B) after multiple inﬁltrations with vanadia, lithiation, and cycling. (D) SEM image of uncoated 3DOM carbon after multiple inﬁltrations with vanadia. Reprinted with permission from .
are inserted in manganese oxides by reduction with LiI in acetonitrile. Finally, 1D nanorods of LiMn2 O4 materials are crystallized into a spinel structure along the [2 2 0] direction by heating under air. In the case of 1D Li[Ni0.25 Li0.15 Mn0.6 ]O2 nanowires with a layered structure (space group: R-3m), they are obtained by one step hydrothermal synthesis without templates . The structure and morphology of the obtained powders are dependent upon both pH and heating temperature. At pH = 10, the spinel phase (space group: Fd-3m) and layered phase (space group: R-3m) were obtained with heating temperature of 150 and 200 ◦ C, respectively. The powder with a layered phase has a nanoplate shape, but was changed into nanowires by controlling the pH from 10 to 2. The Li[Ni0.25 Li0.15 Mn0.6 ]O2 nanowires exhibited good rate performances due to their short diffusion path (95% capacity retention even at 7C rate).
The nanosize effect on rate performance can be maximized by the conﬁguration of the three-dimensionally interpenetrating electrochemical cell structure in which interpenetrating electrode materials are electrically isolated by nanosized ionic conductive layers as an electrolyte. This system was successfully realized by using a photonic crystal structure of three-dimensionally ordered macroporous (3DOM) materials with multiple surface modiﬁcation steps (Fig. 4) . 3DOM materials are inverse replicas of their colloidalcrystal templates, possessing solid nanostructures where voids between closed-packed template spheres once existed. Monolithic three dimensionally ordered macroporous (3DOM) carbon was coated with a conformal layer of poly(phenylene oxide) (PPO) by electropolymerization. The remaining void space was then inﬁltrated with a vanadium alkoxide gel to form interpenetrated vanadium oxide. The
Nanosize in lithium reactive nanomaterials 3DOM carbon and inﬁltrated vanadium oxide have a wall thickness of tens of nm, and the two materials are separated by a few nm PPO polymer layer. Therefore, both the diffusion length in the electrode materials and the conduction length in the electrolytes are on a scale of nm, resulting in decreased mass transport resistances in both the electrode and the electrolyte. 3DOM electrode materials can also be used as composite electrodes of conventional electrochemical cell systems. 3DOM LiFePO4 , LiCoO2 , SnO2 , Si , Li4 Ti5 O12 , V2 O5 , carbon , and LiMn2 O4 spinel  were synthesized using PMMA, PS or silica spheres as closed-packed templates, and exhibited enhanced rate performances. Although nanosized materials clearly show better rate performance due to shorter diffusion length than bulkier materials, from a practical view point some challenges remain. As nanosized materials have higher surface area, more electrolyte is decomposed on the surface of the electrodes, leading to the formation of solid-electrolyte interfaces. This causes poor coulombic efﬁciency, one of the critical parameters for practical utilization of electrode materials. Also, the high surface area could induce poor crystallinity, resulting in fewer storage sites of Li+ ions. In the case of nanosized LiCoO2 , it was found that the coulombic efﬁciencies and reversible capacities decreased to ca. 26% and 65 mAh/g as the size of LiCoO2 is decreased from bulk to 6 nm . Bulkier LiCoO2 usually exhibits >95% and >140 mAh/g, respectively. Therefore, greater focus should be placed on how to obtain more stable surfaces and better crystallinity of nanomaterials for their practical utilization in batteries.
Mechanical effect Energy density as well as safety  and power density are important issues in lithium secondary batteries for EVs. By developing batteries with high energy density electric devices can be operated by batteries for a longer time. Energy density (Wh/g or Wh/cm3 ) is the product of voltage (V) and speciﬁc capacity (Ah/g or Ah/cm3 ). Therefore, it is necessary to increase operating voltage or speciﬁc capacity in order to obtain higher energy density. Conventional strategies for increasing operating voltage and speciﬁc capacity include the development of high voltage cathode materials (>ca. 5 V vs. Li/Li+ ) and high capacity anode materials (>ca. 1000 mAh/g). Changing an electrode material in practical batteries is not simple, however, because it affects the performances of the counter electrode and electrolyte, thus necessitating that all components in the previous system be changed. Above all, although the full cell system is not considered, the successful development of new electrode materials is not straightforward in this regard, because high energy means the material is highly unstable. Alloy electrode materials such as Sn and Si, while presenting inherent challenges, are considered to hold promise for commercialization. The theoretical capacities of Sn and Si materials are 990 and 4200 mAh/g, respectively, much higher than that (370 mAh/g) of graphite. Alloy materials expand and contract as lithium ions are inserted and de-inserted, respectively. The repetition of this volume change during charge and discharge induces fatigue of the
33 electrode materials, and ﬁnally they are degraded by cracking or pulverization, resulting in poor cycle performance [53—56]. As a solution against volume change of alloy materials, nanosize materials have been suggested. Yang et al. reported that nanosize Sn electrodes exhibit much better cycling performance than bulkier Sn electrodes due to a smaller absolute volume change [57,58]. This ﬁnding triggered research on nanostructured alloy materials in lithium ion batteries. The Cho group explored the critical sizes of alloy materials in terms of exhibiting stable cycle performance by the mechanical size effect [59,60]. SnO2 and Si nanomaterials showed the best cycle performances, as their particle size decreased to 3 and 10 nm, respectively. In the case of SnO2 , the smallest nanoparticles showed the best cycle performance. However, Si showed different behavior, and the optimum size was found to be 10 nm. Si nanoparticles with a diameter of 5 nm and 20 nm exhibited poorer electrochemical performances than 10 nm-sized Si. Recently, 1D nanowires of alloy materials have received more attention from researchers than 0D nanoparticles. The Cui group reported that Si and Ge nanowires grown directly on a current collector exhibited stable cycle performances, because the nanowires were not pulverized or broken due to facile strain relaxation of the nanowire geometry (Fig. 5) [61,62]. They showed that the diameter of the nanowires increased without pulverization after several charge-discharge cycles. Si and Ge nanowires were synthesized using a VLS (vapor—liquid—solid) process and a CVD process with SiH4 and GeH4 precursors on stainless steel substrates with Au catalysts. They also synthesized crystalline-amorphous Si core-shell nanowires  and carbon—Si core-shell nanowires  by slight modiﬁcation of the CVD process. In addition to Si nanowires, Sn78 Ge22 —carbon core-shell nanowires were obtained by the Cho group . They ﬁrst synthesized butyl-capped Sn78 Ge22 nanoparticles, and then the nanoparticles were connected to form nanowires by an annealing process. 3D porous structures as well as 1D nanowires have also been utilized in alloy anode materials. The Cho group synthesized mesoporous Si—C core-shell nanowires (Fig. 6) , Si nanotubes , and three-dimensional porous Si particles  using butyl-capped Si solutions with ordered mesoporous silica (SBA-15), a porous anodized alumina membrane, and spherical silica particles as templates, respectively. After cycling, the morphologies were not changed except for an increase in wall thickness. The nanowires delivered speciﬁc capacities of ca. 3000 mAh/g and excellent cycle performances (ca. 90% capacity retention after 100 cycles), as shown in Fig. 7. 3D Si with an inverse opal structure was also synthesized using a CVD method with silica opal templates by Ozin group . Recently, a porous Si—Carbon composite was obtained using a hierarchical bottom-up approach . Si nanoparticles were coated on annealed carbon blacks using a Si CVD process, and then Si-coated carbon black particles were assembled into rigid spherical granules using a carbon CVD process. In addition to Si, 3D porous structures have been used in various alloy electrode materials. The Cho group synthesized ordered mesoporous Sn2 P2 O7 , 3D porous Ge , and hollow Sb nanoparticles  using a surfactant, silica spheres, and silica spheres as templates, respectively.
K.T. Lee, J. Cho
Figure 5 Morphology and electronic changes in Si NWs from reaction with Li. SEM image of pristine Si NWs before (a) and after (b) electrochemical cycling (same magniﬁcation). The inset in (a) is a cross-sectional image showing that the NWs are directly contacting the stainless steel current collector. TEM image of a pristine Si NW with a partial Ni coating before (c) and after (d) Li cycling. (e) Size distribution of NWs before and after charging to 10 mV (bin width 10 nm). The average diameter of the NWs increased from 89 to 141 nm after lithiation. (f) I—V curve for a single NW device (SEM image, inset) constructed from a pristine Si NW. (g) I—V curve for a single NW device (SEM image, inset) constructed from a NW that had been charged and discharged once at a C/20 rate. Reprinted with permission from [61,62].
The electrochemical performances of nanosized Sn are more enhanced in combination with carbon nanostructure. In 2003, Lee et al. suggested the use of a cavity structure to buffer the volume change of Sn by synthesis of nano Sn-encapsulated hollow carbon composite (Fig. 8) . This synthetic method exploits the different hydrophobicity between the Sn precursor (Tert-butyl phenyl tin) and carbon precursor (Resorcinol-Formaldehyde). By sol—gel polymerization of resorcinol-formaldehyde (RF) with tributyl phenyl tin (TBPT), it forms a core-shell structure of TBPT-RF. A cavity structure is then generated by decomposition of an organometallic Sn precursor into Sn metal during carbonization of the RF polymer. The disadvantage of this method is that TBPT has a boiling point in a range of 125—128 ◦ C @
0.14 mmHg, and thus the decomposition and vaporization of TBPT are competitive during the carbonization reaction. Therefore, some TBPT can evaporate when the RF shell is too thin, resulting in a limitation of maximizing Sn content in the composite. The tin content in the composite was ca. 24 wt%, and it delivered 452 mAh/g. As a reference, Sn-free carbon delivered 373 mAh/g. Independently, this synthetic method was recently revisited by the Scrosati group in 2007 and 2008 [74,75]. They used the same precursors of TBPT and RF chemicals, and slightly modiﬁed the synthetic method including variation of the HCHO concentration and the sequence of mixing precursors. The tin content in the Sn/C composite was increased to ca. 50 wt%, but only ca. 500 mAh/g was delivered. The small
Nanosize in lithium reactive nanomaterials
Figure 6 (a) Schematic illustration of the preparation of [email protected]
core-shell nanowires. (b) TEM image of [email protected]
core-shell nanorods obtained from ﬁrst impregnation. (c) Expanded TEM image of (b) (inset is SADP of (c)). (d) TEM image of [email protected]
coreshell nanowires obtained from fourth impregnation. (e) Expanded TEM image of (d) and (f) Raman spectrum of [email protected]
core-shell nanowires. Reprinted with permission from [31,32].
speciﬁc capacity of the Sn/C composite was attributed to the much smaller speciﬁc capacity of the RF carbon material (45 mAh/g) relative to that of conventional RF carbons (>300 mAh/g). This might be caused by low carbonization temperature, 700 ◦ C. More tin-encapsulated hollow carbon materials were synthesized, and synthetic methods were amended to increase the loading amount of Sn nanoparticles, resulting in higher speciﬁc capacities. Allyltriphenyltin was simply heated through several temperature steps in a sealed vacuum tube, and it forms a hollow tin—carbon core-shell structure . Hollow tin had an irregular wall structure with a thickness of 2—6 nm, and the thickness of the carbon layer was 2—3 nm. The tin content in the composite was ca. 43 wt% and its reversible capacity was 550 mAh/g. A hard template
method was also adopted to obtain tin-encapsulated hollow carbons . SnO2 was coated on silica spheres, and the carbon precursor was coated on SnO2 after dissolution of the silica template. Sn-encapsulated hollow carbon was then obtained by heating at 700 ◦ C. The Sn content was increased to 70 wt% by controlling the thickness of SnO2 and carbon layers, resulting in a high speciﬁc capacity of >800 mAh/g. Similar cavity approaches were utilized using various carbon structures such as carbon nanotubes, microtubes, and hollow nanoﬁbers. A Sn/CNT composite was synthesized using AAO templates . Liquid SnCl4 or 1 M SnCl4 solution was inﬁltrated into the AAO template, and carbon was coated on the SnCl4 /AAO composites by the CVD method with C2 H2 , accompanying a reduction of SnCl4 to Sn metal. Also, Sn-encapsulated microtubes  or hollow
K.T. Lee, J. Cho
Figure 7 (a) Voltage proﬁles of a [email protected]
core-shell nanowire electrode after 1, 30, 60, and 80 cycles at a rate of 0.2C between 1.5 and 0 V in coin-type half-cells. (b) Plot of charge capacity and Coulombic efﬁciency of the cell (a) vs. cycle number. (c) Differential curves of the [email protected]
core-shell nanowire electrode after 1, 2, and 30 cycles. (d) Voltage proﬁles of the rate capabilities of the [email protected]
core-shell nanowire electrode at rates of 0.2, 0.5, 1, and 2C between 1.5 and 0 V in coin-type half-cells (same charge and discharge rates were used). Reprinted with permission from .
nanoﬁbers  were obtained by an electrospinning process using a single-needle nozzle. Viscous liquid tin and carbon precursors (organometallic tin and poly-acrylonitrile) were electrospun to form microtubes or nanoﬁbers, and they were subsequently carbonized to form Sn-encapsulated hollow carbon microtubes or nanoﬁbers. Similar morphology was utilized in a Si anode material. CNT was obtained by a CVD process using alumina tube templates. Si was then coated
by SiH4 decomposition followed by dissolution of alumina templates in an HF solution .
Structural effect The above two roles of nanomaterials were easily anticipated, but unexpected phenomena of nanomaterials have
Figure 8 Transmission electron micrograph of the core-she6ll structure of tributylphenyltin/RF polymer (a) and Sn-encapsulated spherical hollow carbon. (b) Field-emission scanning electron micrograph of TC1 (c) and TC2 (d). (e) Synthetic scheme for Snencapsulated spherical hollow carbon. Reprinted with permission from .
Nanosize in lithium reactive nanomaterials
Figure 9 (a) Schematic derivation of the size-dependent FePO4 —LiFePO4 binary phase diagram using the XRD data for Li0.60 FePO4 . (b) Open-circuit voltage vs. x in Lix FePO4 at 25 ◦ C. Voltage was measured after discharging (starting from FePO4 ) or charging (starting from LiFePO4 ) a cell at 150 A to each measuring point, followed by a wait of 30 h to reach equilibrium. Reprinted with permission from .
recently been found. Maier et al. suggested that the redox potential decreases due to increased interface energy as the size of the electrode materials decreases [82—84]. However, the Chiang group showed an opposite phenomenon for LiFePO4 through careful experiments, although no clear explanation for this ﬁnding was given . Besides the redox potential, LiFePO4 exhibits various interesting characteristics related to nanosize. As the size of LiFePO4 decreases, it is found that the miscibility gap decreases (Fig. 9) . The miscibility gap of LiFePO4 indicates a region of coexistence of Li-poor (Li˛ FePO4 ) and Li-rich phases (Li1−␤ FePO4 ). This means that solid solution behavior is exhibited in a range of x ≤ ˛ and x ≥ 1 − ˇ for x in Lix FePO4 . For bulk LiFePO4 materials, the experimental values of ˛ and ˇ are 0.05 and 0.11, respectively , and these values increased to 0.12 and 0.17 as the particle size was decreased to 34 nm, as conﬁrmed by electrochemical measurements  and XRD experiments  (Fig. 10). This phenomenon is attributed to the coherence stress. As the size of LiFePO4 decreases, the stress within two phases of Li˛ FePO4 and Li1−ˇ FePO4 becomes signiﬁcant. This makes the coexistence of two phases in a single particle unstable, resulting in a decrease of the miscibility gap. The similar phenomenon was observed in anatase TiO2 nanoparticles. Wagemaker et al. investi-
Figure 10 Left: magniﬁed view of the 200 peak measured for Li0.93 FePO4 with different particle sizes of 200, 80, and 40 nm at room temperature. Right: derivation of changes in the a-axis lattice parameter of lithium-rich tryphilite phase as a function of lithium composition x for different particle sizes. Reprinted with permission from .
gated the phase diagram of the Li composition in anatase TiO2 versus crystal particle size based on the neutron diffraction results . The phase diagram changes as the particle size decreases. Decreasing particle size causes to increase the solid solution region of Li˛ TiO2 (˛ = 0.03 and 0.21 for 120 nm and 7 nm TiO2 , respectively) and form new Li-rich phase (Li1 TiO2 ) not observed in bulk TiO2 materials. In the case of 7 nm anatase TiO2 , the reaction of Li insertion proceeds through three steps: one phase reaction of anatase phase (I41/amd)—–two phase reaction of anatase and Lititanate phase (Imma)—–two phase reaction of Li—titanate phase (Imma) and Li1 TiO2 (I41/amd). However, Bruce group recently showed that the ordered mesoporous anatase TiO2 (6.5 nm wall thickness) exhibited a different reaction mechanism from the 7 nm anatase TiO2 nanoparticles: one phase reaction of anatase phase up to Li0.05 TiO2 —–two phase
K.T. Lee, J. Cho
Figure 11 (a) Raman spectra taken on the LiGa phase derived from Ga and CuGa2 electrodes. (b) Ex situ XRD patterns of LiGa phase. The samples for Raman and ex situ XRD analyses were prepared by lithiating the CuGa2 and Ga electrode down to 0.0 V (vs. Li/Li+ ) (temperature = 25 ◦ C and speciﬁc current = 10 mA/gGa ). (c) d-spacing values of LiGa phase derived from the CuGa2 electrode (temperature = 120 ◦ C). The triangles, circles, and squares denote d111 , d220 , and d311 , respectively. (d) Delithiation rate performance of Ga and CuGa2 electrodes (0.0—2.0 V vs. Li/Li+ ). The electrodes were lithiated up to the LiGa phase to compare their delithiation rate performance from LiGa to pure Ga. Cycle performance is also compared for the Ga and CuGa2 electrode in the inset. Reprinted with permission from .
reaction of anatase phase and Li0.45 TiO2 —–one phase reaction of Li0.45 TiO2 (continuous lattice expansion up to Li0.96 TiO2 ) . Also, the surface area of LiFePO4 increases as the particle size decreases, causing sufﬁcient surface oxidation to change the lattice parameters of LiFePO4 nanoparticles . Fe2+ on the surface of LiFePO4 is changed to Fe3+ due to surface oxidation upon exposure to air, leading to the formation of a Li1−x FePO4 solid solution. Therefore, the volume of LiFePO4 nanoparticles becomes lower than that of bulkier LiFePO4 , when they are exposed to air. Another interesting nano-size effect was found in transition metal oxide anode materials. As transition metal oxides such as CoO, Co3 O4 , CuO, NiO, and MnO are fully reduced by electrochemical insertion of lithium ions (MOx + 2xLi ↔ M + xLi2 O), the Tarascon group reported that these oxides formed composite materials comprised of nanosized metallic clusters (2—8 nm) dispersed in an amorphous Li2 O matrix . The conversion reaction of transition metal oxides with lithium was found to be electrochemically reversible for hundreds of cycles [91—93]. It was previously thought that this conversion reaction is irreversible and only electrochemical reduction of the metal oxides
occurs, although an exceptional case of molybdenum oxides had been previously introduced by the Nazar group . However, the large contact area between nanosized metallic clusters and the Li2 O matrix in the reduced composites enables the reverse oxidation reaction electrochemically. The transition metal oxides have an advantage of high speciﬁc capacities (>ca. 700 mAh/g); however, they also have a drawback of high polarization due to poor kinetics of conversion reactions, which is attributed to the high barrier energy to activate the breaking of the M—O bonds. The same drawback was also applied to transition metal phosphide anode materials such as FeP2 , ZnP2 , and NiP2 , although they have lower polarization than oxide materials [95—97]. A similar phenomenon was found in binary alloy anode materials. Lee et al. found that the conversion reaction of binary intermetallics such as CuGa2 and NiGa4 forms alloy composite materials composed of nanosized inactive metallic (Cu, Ni) grains and active Li—Ga alloy grains . The electrochemical behavior of (Li—Ga + M) nanocomposites obtained from binary intermetallics is similar to that of (Li2 O + M) nanocomposites obtained from metal oxides. Both exhibited high polarization caused by high activation
Nanosize in lithium reactive nanomaterials energy to break the M—Ga and M—O bonds. Therefore, the overpotential of lithiation reactions was much higher than that of delithiation reactions. They also further examined the role of electrochemically driven inactive metallic grains in the electrochemical performances of binary intermetallics. They proposed that there is a partial bonding between two nanograins, based on Raman and XRD analyses including Raman peak shift and the different lattice parameters of the electrochemically driven LiGa phase from the binary intermetallics (Fig. 11a and b). This interfacial phenomenon resulted in unexpected thermodynamic and kinetic properties. The nanocomposites of Li2 Ga + M showed a different delithiation pathway including different solid solution behavior and the formation of a new Li—Ga phase (Fig. 11c), which was not seen in Li2 Ga obtained from pure Ga [98,99]. Above all, binary Ga—M intermetallics showed exceptionally good delithiation rate performances, better than pure Ga, which is attributed to the weakened Li—Ga bond as a result of M—Ga partial bonding (Fig. 11d). Similar behaviors were also found in In—Cu intermetallics , MoO2 , and MGeO3 (M = Cu, Fe, Co)  materials. It is considered that this mechanism can be further utilized to design new materials for high rate electrode materials.
Summary and outlook In this review, the roles of nanosize in lithium reactive nanomaterials in terms of the electrochemical performances of lithium ion batteries were highlighted. The following main points features were discussed: (i) Nanostructured materials show excellent rate performances due to short diffusion length. Various morphologies of nanostructured materials such as 1D nanorods, 2D plates, and 3D porous materials have been synthesized through various synthetic routes. (ii) Nanostructured materials exhibited peculiar mechanical properties. In the case of alloy anode materials, they have a fatal problem of pulverization and cracking of electrode materials by volume change during charging and discharging, resulting in poor cycle performance. This problem can be solved by utilizing 1D nanowire materials or alloy-encapsulated hollow structures due to their stable mechanical structure. (iii) Nanostructured materials also showed unexpected and interesting phenomena, referred to as structural effects. Nanosize in electrode materials affects not only thermodynamic but also kinetic properties. It changed the electrochemical redox potential of LiFePO4 cathode materials and also enhanced the rate capability of binary intermetallic anode materials. These roles of nanosize in electrode materials have been emphasized in conventional batteries. However, conventional batteries based on intercalation chemistry have a limitation with respect to increasing the battery energy density, due to topotactic redox reactions. Their maximum energy densities are not adequate to meet the demands of new markets such as electric vehicles. Therefore, new electrochemical systems with higher energy density are being sought for next-generation batteries, and Li—air and Li—sulfur batteries are considered promising alternatives. Their theoretical energy densities are more than 2500 Wh/kg, a factor of at least 3—5 times higher relative to conventional batteries.
39 It is believed that nanostructured materials will also prove highly useful in Li—air and Li—sulfur batteries. Recently, ordered mesoporous carbon was used as an effective sulfur reservoir for Li—sulfur batteries [103—105]. The sulfur conﬁned in a mesoporous carbon composite improved electrical conductivity due to essential contact between the conductive carbon framework and sulfur, and it also showed a kinetic inhibition effect by trapping dissolved polysulﬁdes within the carbon framework . The nanostructure in the composite, therefore, resulted in improved battery performance. Also, nanostructured metal oxides were recently used as catalysts in Li-air cells [106—108]. Delithiation reaction of lithium oxides was previously thought to be an irreversible reaction. However, nanostructured materials showed excellent catalytic properties, exhibiting a reversible redox reaction between lithium and oxygen. However, nanostructured materials have thus far not been extensively studied in Li—air and Li—sulfur batteries. Therefore, additional research to explore their potential is required in order to overcome the problems restricting application of these new systems.
Acknowledgements This research was supported by the Converging Research Center Program through the Ministry of Education, Science and Technology (2010K000984). Also, Support from the National Research Foundation of Korea (NRF) grant funded by the Korea Government (MEST) (No. 2010-0019408) is greatly acknowledged.
References  P.G. Bruce, B. Scrosati, J.M. Tarascon, Angewandte ChemieInternational Edition 47 (2008) 2930—2946.  B.L. Ellis, K.T. Lee, L.F. Nazar, Chemistry of Materials 22 (2010) 691—714.  P. Gibot, M. Casas-Cabanas, L. Laffont, S. Levasseur, P. Carlach, S. Hamelet, J.M. Tarascon, C. Masquelier, Nature Materials 7 (2008) 741—747.  A.K. Padhi, K.S. Nanjundaswamy, J.B. Goodenough, Journal of the Electrochemical Society 144 (1997) 1188—1194.  G.Y. Chen, X.Y. Song, T.J. Richardson, Electrochemical and Solid State Letters 9 (2006) A295—A298.  L. Laffont, C. Delacourt, P. Gibot, M.Y. Wu, P. Kooyman, C. Masquelier, J.M. Tarascon, Chemistry of Materials 18 (2006) 5520—5529.  C. Delmas, M. Maccario, L. Croguennec, F. Le Cras, F. Weill, Nature Materials 7 (2008) 665—671.  K.T. Lee, W.H. Kan, L.F. Nazar, Journal of the American Chemical Society 131 (2009) 6044.  B. Kang, G. Ceder, Nature 458 (2009) 190—193.  K. Zaghib, J.B. Goodenough, A. Mauger, C. Julien, Journal of Power Sources 194 (2009) 1021—1023.  G. Ceder, B. Kang, Journal of Power Sources 194 (2009) 1024—1028.  C. Delacourt, P. Poizot, S. Levasseur, C. Masquelier, Electrochemical and Solid State Letters 9 (2006) A352—A355.  Y. Ren, A.R. Armstrong, F. Jiao, P.G. Bruce, Journal of the American Chemical Society 132 (2010) 996—1004.  K. Amine, I. Belharouak, Z.H. Chen, T. Tran, H. Yumoto, N. Ota, S.T. Myung, Y.K. Sun, Advanced Materials 22 (2010) 3052—3057.
40  D. Morgan, A. Van der Ven, G. Ceder, Electrochemical and Solid State Letters 7 (2004) A30—A32.  M.S. Islam, D.J. Driscoll, C.A.J. Fisher, P.R. Slater, Chemistry of Materials 17 (2005) 5085—5092.  S. Nishimura, G. Kobayashi, K. Ohoyama, R. Kanno, M. Yashima, A. Yamada, Nature Materials 7 (2008) 707—711.  J.Y. Li, W.L. Yao, S. Martin, D. Vaknin, Solid State Ionics 179 (2008) 2016—2019.  R. Amin, J. Maier, P. Balaya, D.P. Chen, C.T. Lin, Solid State Ionics 179 (2008) 1683—1687.  Y.G. Wang, Y.R. Wang, E.J. Hosono, K.X. Wang, H.S. Zhou, Angewandte Chemie-International Edition 47 (2008) 7461—7465.  S.K. Martha, J. Grinblat, O. Haik, E. Zinigrad, T. Drezen, J.H. Miners, I. Exnar, A. Kay, B. Markovsky, D. Aurbach, Angewandte Chemie-International Edition 48 (2009) 8559—8563.  S.Y. Chung, J.T. Bloking, Y.M. Chiang, Nature Materials 1 (2002) 123—128.  D.H. Kim, J. Kim, Electrochemical and Solid State Letters 9 (2006) A439—A442.  A.V. Murugan, T. Muraliganth, A. Manthiram, Electrochemistry Communications 10 (2008) 903—906.  B. Ellis, W.H. Kan, W.R.M. Makahnouk, L.F. Nazar, Journal of Materials Chemistry 17 (2007) 3248—3254.  S.F. Yang, P.Y. Zavalij, M.S. Whittingham, Electrochemistry Communications 3 (2001) 505—508.  D.Y. Wang, H. Buqa, M. Crouzet, G. Deghenghi, T. Drezen, I. Exnar, N.H. Kwon, J.H. Miners, L. Poletto, M. Graetzel, Journal of Power Sources 189 (2009) 624—628.  D.W. Choi, D.H. Wang, I.T. Bae, J. Xiao, Z.M. Nie, W. Wang, V.V. Viswanathan, Y.J. Lee, J.G. Zhang, G.L. Graff, Z.G. Yang, J. Liu, Nano Letters 10 (2010) 2799—2805.  T. Yoshida, K. Kitoh, S. Ohtsubo, W. Shionoya, H. Katsukawa, J. Yamaki, Electrochemical and Solid State Letters 10 (2007) A60—A64.  W. Lu, I. Belharouak, S.H. Park, Y.K. Sun, K. Amine, Electrochimica Acta 52 (2007) 5837—5842.  F. Jiao, J.L. Bao, A.H. Hill, P.G. Bruce, Angewandte ChemieInternational Edition 47 (2008) 9711—9716.  J.Y. Luo, Y.G. Wang, H.M. Xiong, Y.Y. Xia, Chemistry of Materials 19 (2007) 4791—4795.  F. Jiao, K.M. Shaju, P.G. Bruce, Angewandte ChemieInternational Edition 44 (2005) 6550—6553.  S.M. Zhu, H.S. Zhou, T. Miyoshi, M. Hibino, I. Honma, M. Ichihara, Advanced Materials 16 (2004) 2012.  X.X. Li, F.Y. Cheng, B. Guo, J. Chen, Journal of Physical Chemistry B 109 (2005) 14017—14024.  A.L.M. Reddy, M.M. Shaijumon, S.R. Gowda, P.M. Ajayan, Nano Letters 9 (2009) 1002—1006.  M.H. Lee, J.Y. Kim, H.K. Song, Chemical Communications 46 (2010) 6795—6797.  M.M. Ren, Z. Zhou, X.P. Gao, L. Liu, W.X. Peng, Journal of Physical Chemistry C 112 (2008) 13043—13046.  J.Y. Luo, H.M. Xiong, Y.Y. Xia, Journal of Physical Chemistry C 112 (2008) 12051—12057.  D.K. Kim, P. Muralidharan, H.W. Lee, R. Ruffo, Y. Yang, C.K. Chan, H. Peng, R.A. Huggins, Y. Cui, Nano Letters 8 (2008) 3948—3952.  M.G. Kim, M. Jo, Y.S. Hong, J. Cho, Chemical Communications (2009) 218—220.  N.S. Ergang, J.C. Lytle, K.T. Lee, S.M. Oh, W.H. Smyrl, A. Stein, Advanced Materials 18 (2006) 1750.  C.M. Doherty, R.A. Caruso, B.M. Smarsly, C.J. Drummond, Chemistry of Materials 21 (2009) 2895—2903.  N.S. Ergang, J.C. Lytle, H.W. Yan, A. Stein, Journal of the Electrochemical Society 152 (2005) A1989—A1995.  J.C. Lytle, H.W. Yan, N.S. Ergang, W.H. Smyrl, A. Stein, Journal of Materials Chemistry 14 (2004) 1616—1622.
K.T. Lee, J. Cho  A. Esmanski, G.A. Ozin, Advanced Functional Materials 19 (2009) 1999—2010.  E.M. Sorensen, S.J. Barry, H.K. Jung, J.R. Rondinelli, J.T. Vaughey, K.R. Poeppelmeier, Chemistry of Materials 18 (2006) 482—489.  H.W. Yan, S. Sokolov, J.C. Lytle, A. Stein, F. Zhang, W.H. Smyrl, Journal of the Electrochemical Society 150 (2003) A1102—A1107.  K.T. Lee, J.C. Lytle, N.S. Ergang, S.M. Oh, A. Stein, Advanced Functional Materials 15 (2005) 547—556.  D. Tonti, M.J. Torralvo, E. Enciso, I. Sobrados, J. Sanz, Chemistry of Materials 20 (2008) 4783—4790.  M. Okubo, E. Hosono, J. Kim, M. Enomoto, N. Kojima, T. Kudo, H.S. Zhou, I. Honma, Journal of the American Chemical Society 129 (2007) 7444—7452.  Y.D. Wang, J.W. Jiang, J.R. Dahn, Electrochemistry Communications 9 (2007) 2534—2540.  L.Y. Beaulieu, S.D. Beattie, T.D. Hatchard, J.R. Dahn, Journal of the Electrochemical Society 150 (2003) A419—A424.  L.Y. Beaulieu, T.D. Hatchard, A. Bonakdarpour, M.D. Fleischauer, J.R. Dahn, Journal of the Electrochemical Society 150 (2003) A1457—A1464.  J.H. Ryu, J.W. Kim, Y.E. Sung, S.M. Oh, Electrochemical and Solid State Letters 7 (2004) A306—A309.  M.M. Thackeray, J.T. Vaughey, C.S. Johnson, A.J. Kropf, R. Benedek, L.M.L. Fransson, K. Edstrom, Journal of Power Sources 113 (2003) 124—130.  J. Yang, M. Winter, J.O. Besenhard, Solid State Ionics 90 (1996) 281—287.  J. Yang, M. Wachtler, M. Winter, J.O. Besenhard, Electrochemical and Solid State Letters 2 (1999) 161—163.  C. Kim, M. Noh, M. Choi, J. Cho, B. Park, Chemistry of Materials 17 (2005) 3297—3301.  H. Kim, M. Seo, M.H. Park, J. Cho, Angewandte ChemieInternational Edition 49 (2010) 2146—2149.  C.K. Chan, H.L. Peng, G. Liu, K. McIlwrath, X.F. Zhang, R.A. Huggins, Y. Cui, Nature Nanotechnology 3 (2008) 31—35.  C.K. Chan, X.F. Zhang, Y. Cui, Nano Letters 8 (2008) 307—309.  L.F. Cui, R. Ruffo, C.K. Chan, H.L. Peng, Y. Cui, Nano Letters 9 (2009) 491—495.  L.F. Cui, Y. Yang, C.M. Hsu, Y. Cui, Nano Letters 9 (2009) 3370—3374.  H. Lee, J. Cho, Nano Letters 7 (2007) 2638—2641.  H. Kim, J. Cho, Nano Letters 8 (2008) 3688—3691.  M.H. Park, M.G. Kim, J. Joo, K. Kim, J. Kim, S. Ahn, Y. Cui, J. Cho, Nano Letters 9 (2009) 3844—3847.  H. Kim, B. Han, J. Choo, J. Cho, Angewandte ChemieInternational Edition 47 (2008) 10151—10154.  A. Magasinski, P. Dixon, B. Hertzberg, A. Kvit, J. Ayala, G. Yushin, Nature Materials 9 (2010) 353—358.  E. Kim, D. Son, T.G. Kim, J. Cho, B. Park, K.S. Ryu, S.H. Chang, Angewandte Chemie-International Edition 43 (2004) 5987—5990.  M.H. Park, K. Kim, J. Kim, J. Cho, Advanced Materials 22 (2010) 415.  H. Kim, J. Cho, Chemistry of Materials 20 (2008) 1679—1681.  K.T. Lee, Y.S. Jung, S.M. Oh, Journal of the American Chemical Society 125 (2003) 5652—5653.  J. Hassoun, G. Derrien, S. Panero, B. Scrosati, Advanced Materials 20 (2008) 3169—3175.  G. Derrien, J. Hassoun, S. Panero, B. Scrosati, Advanced Materials 19 (2007) 2336.  G.L. Cui, Y.S. Hu, L.J. Zhi, D.Q. Wu, I. Lieberwirth, J. Maier, K. Mullen, Small 3 (2007) 2066—2069.  W.M. Zhang, J.S. Hu, Y.G. Guo, S.F. Zheng, L.S. Zhong, W.G. Song, L.J. Wan, Advanced Materials 20 (2008) 1160.  Y. Wang, M. Wu, Z. Jiao, J.Y. Lee, Chemistry of Materials 21 (2009) 3210—3215.
Nanosize in lithium reactive nanomaterials  Y. Yu, L. Gu, C.B. Zhu, P.A. van Aken, J. Maier, Journal of the American Chemical Society 131 (2009) 15984.  Y. Yu, L. Gu, C.L. Wang, A. Dhanabalan, P.A. van Aken, J. Maier, Angewandte Chemie-International Edition 48 (2009) 6485—6489.  B. Hertzberg, A. Alexeev, G. Yushin, Journal of the American Chemical Society 132 (2010) 8548.  J. Maier, Nature Materials 4 (2005) 805—815.  J. Maier, Journal of Power Sources 174 (2007) 569—574.  J. Jamnik, J. Maier, Physical Chemistry Chemical Physics 5 (2003) 5215—5220.  N. Meethong, H.Y.S. Huang, W.C. Carter, Y.M. Chiang, Electrochemical and Solid State Letters 10 (2007) A134—A138.  G. Kobayashi, S.I. Nishimura, M.S. Park, R. Kanno, M. Yashima, T. Ida, A. Yamada, Advanced Functional Materials 19 (2009) 395—403.  A. Yamada, H. Koizumi, S.I. Nishimura, N. Sonoyama, R. Kanno, M. Yonemura, T. Nakamura, Y. Kobayashi, Nature Materials 5 (2006) 357—360.  M. Wagemaker, W.J.H. Borghols, F.M. Mulder, Journal of the American Chemical Society 129 (2007) 4323—4327.  Y. Ren, L.J. Hardwick, P.G. Bruce, Angewandte ChemieInternational Edition 49 (2010) 2570—2574.  P. Poizot, S. Laruelle, S. Grugeon, L. Dupont, J.M. Tarascon, Nature 407 (2000) 496—499.  F. Badway, I. Plitz, S. Grugeon, S. Laruelle, M. Dolle, A.S. Gozdz, J.M. Tarascon, Electrochemical and Solid State Letters 5 (2002) A115—A118.  D. Larcher, G. Sudant, J.B. Leriche, Y. Chabre, J.M. Tarascon, Journal of the Electrochemical Society 149 (2002) A234—A241.  H. Li, P. Balaya, J. Maier, Journal of the Electrochemical Society 151 (2004) A1878—A1885.  F. Leroux, L.F. Nazar, Solid State Ionics 133 (2000) 37—50.  F. Gillot, S. Boyanov, L. Dupont, M.L. Doublet, A. Morcrette, L. Monconduit, J.M. Tarascon, Chemistry of Materials 17 (2005) 6327—6337.  D.C.C. Silva, O. Crosnier, G. Ouvrard, J. Greedan, A. SafaSefat, L.F. Nazar, Electrochemical and Solid State Letters 6 (2003) A162—A165.  H. Hwang, M.G. Kim, Y. Kim, S.W. Martin, J. Cho, Journal of Materials Chemistry 17 (2007) 3161—3166.  K.T. Lee, Y.S. Jung, J.Y. Kwon, J.H. Kim, S.M. Oh, Chemistry of Materials 20 (2008) 447—453.  K.T. Lee, Y.S. Jung, T. Kim, C.H. Kim, J.H. Kim, J.Y. Kwon, S.M. Oh, Electrochemical and Solid State Letters 11 (2008) A21—A24.
41  Y.S. Jung, K.T. Lee, J.H. Kim, J.Y. Kwon, S.M. Oh, Advanced Functional Materials 18 (2008) 3010—3017.  J.H. Ku, Y.S. Jung, K.T. Lee, C.H. Kim, S.M. Oh, Journal of the Electrochemical Society 156 (2009) A688— A693.  C.H. Kim, Y.S. Jung, K.T. Lee, J.H. Ku, S.M. Oh, Electrochimica Acta 54 (2009) 4371—4377.  X.L. Ji, K.T. Lee, L.F. Nazar, Nature Materials 8 (2009) 500—506.  C.D. Liang, N.J. Dudney, J.Y. Howe, Chemistry of Materials 21 (2009) 4724—4730.  Y. Yang, M.T. McDowell, A. Jackson, J.J. Cha, S.S. Hong, Y. Cui, Nano Letters 10 (2010) 1486—1491.  T. Ogasawara, A. Debart, M. Holzapfel, P. Novak, P.G. Bruce, Journal of the American Chemical Society 128 (2006) 1390—1393.  A. Debart, A.J. Paterson, J. Bao, P.G. Bruce, Angewandte Chemie-International Edition 47 (2008) 4521— 4524.  Y.C. Lu, Z.C. Xu, H.A. Gasteiger, S. Chen, K. Hamad-Schifferli, Y. Shao-Horn, Journal of the American Chemical Society 132 (2010) 12170—12171. Kyu Tae Lee is a Professor of the Interdisciplinary School of Green Energy at Ulsan National Institute of Science and Technology (UNIST), Korea. He received his B.Sc. (2000) and Ph.D. (2006) in Chemical Engineering from Seoul National University (Korea). In 2007, he joined Prof. Linda F. Nazar group at University of Waterloo (Canada) as a post-doctoral fellow. His current research is focused on nanostructured materials for Li-ion and Na-ion batteries and new electrochemical energy conversion/storage systems. Jaephil Cho is a Professor and a dean in Interdisciplinary School of Green Energy at UNIST (Korea). He received his Ph.D. (1995) in Ceramic Engineering from Iowa State University (USA). Currently, he is a director of Converging Research Center for Innovative Battery Technologies funded by Ministry of Education, Science and Technology in Korea. His current research is focused mainly on nanomaterials for energy conversion and storage, nanoscale coating, and safety enhancement of the Li-ion batteries.