SECONDARY BATTERIES – LITHIUM RECHARGEABLE SYSTEMS – LITHIUM-ION | Electrolytes: Solid Oxide

Electrolytes: Solid Oxide K Takada, National Institute for Materials Science, Tsukuba, Japan & 2009 Elsevier B.V. All rights reserved.

Introduction Today, lithium-ion batteries are a key component in advanced information society, which are used to power portable information equipment including mobile phones and notebook PCs. In addition, they are required to contribute to the realization of environment-friendly society as the power source for electric vehicles and loadleveling apparatus. However, safety issue is intrinsic for the lithium-ion batteries, because lithium-ion batteries contain combustible substances, which are organic solvents used in the electrolytes. Solid electrolytes are expected to be the fundamental solution for the safety issue, and all-ceramic solid-state batteries are strongly longed for to satisfy the global needs. Another advantage of solid-state batteries is that they show long life. Every species other than Liþ ions, for example, counter-anions, solvent molecules, and impurities, migrate in liquid electrolyte. When such a species migrates to the surface of cathode, it may be oxidatively decomposed; or one reaching the surface of the anode may be reduced. That is, the migration can cause side reactions in the batteries to degrade the battery performance. On the contrary, only lithium ions diffuse in solid electrolytes. It means that there is no movement of such species making side reactions proceed. Therefore, solid-state batteries show long cycle life, long shelf life, and extremely small self-discharge. Highly conductive solid electrolyte is a key in the development of the solid-state battery. There are several types of solid electrolytes: nitrides, halides, sulfides, and oxides, among which oxides have the following advantages. The first advantage is their chemical stability. Other electrolytes are, for example, hygroscopic, and difficult to handle in ambient atmosphere, which is a drawback in assembling batteries or components. One can employ a variety of processing methods in the assembly because of the stability. Electrolyte layers should be as thin as possible in order to reduce the internal resistances of the batteries. Such thin electrolyte layers are, in general, fabricated by thin-film technology, including evaporation process. However, such process may not fit mass production of batteries. On the contrary, oxide solid electrolytes can be formed into thin films by wet process, including sol–gel method, which will be much better for the battery production. Finally, oxide ions are high in electronegativity. It means that electrons

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are hardly removed from the oxide electrolytes, and they are stable against electrochemical oxidation. Therefore, the oxide electrolytes will be compatible with highvoltage cathode. In general, oxides are relatively easy to be synthesized, and thus several solid electrolytes have been found. The first oxide electrolyte with lithium-ionic conduction was Li-b-alumina, and the finding was followed by many studies aiming at the development of highly conductive oxide electrolytes. Around 1980, most of the studies were on lithium-ionic conduction in lithium-rich oxysalts, and some oxysalts, many of which have g-Li3PO4 structure, were developed. Because their ionic conductivities are of the orders from 107 to 105 S cm1 at room temperature, which are lower than that in liquids by several orders, they were applied to thin-film batteries in order to make up for the poor ionic conduction and make the internal resistance as small as possible. Drastic increase of the conductivity in oxides was achieved in the 1990s. Two types of oxide electrolytes were developed: one has a sodium super ionic conductor (NASICON)-type structure, and the other has a perovskite-type structure. Both of them show ionic conductivities of 103 S cm1. Details of the three types of oxide solid electrolytes – g-Li3PO4 type, NASICON type, and perovskite type – are described in this article.

c-Li3PO4-Type Oxysalts Some oxysalts including Li2SO4 show fast ionic conduction in their high-temperature phases, and many attempts were made to stabilize such highly conducting phases down to ambient temperature. In 1977, R. D. Shannon and coworkers synthesized Li2þxC1xBxO3, Li3xB1xCxO3, Li4xSi1xPxO4, Li42xSi1xSxO4, Li4þxSi1xAlxO4, and Li5xAl1xSixO4 and reported that ionic conduction is enhanced when two types of oxysalts are formed into solid solutions. The study was followed by a finding of fast ionic conduction in Li14Zn(GeO4)4 in 1978. The material was named lithium (LISICON) (lithium superionic conductor), motivated such studies, and many lithium-ion-conducting oxysalts were found around 1980: one group comprises solid solutions based on lithium nesosilicate (Li4SiO4), and the other is based on g-Li3PO4. Besides R. D. Shannon’s work, systematic study was also done by A. R. West and coworkers on xLi4MIVO4– (1 – x)Li3MVO4 systems (MIV ¼ Ge, Ti and MV ¼As, V).

Secondary Batteries – Lithium Rechargeable Systems – Lithium-Ion | Electrolytes: Solid Oxide

The conductivities of the solid solutions always showed maxima at x ¼ 0.4–0.6, which was consistent with R. D. Shannon’s results. On the contrary, comparison between the ionic conductivities of different types of oxysalts revealed strong correlation between the unit cell volume and the conductivity. Figure 1 shows the relationship between the cell volume and the ion-conducting properties of the eight systems. The conductivity increased and the activation energy for conduction decreased with increasing cell volume, while the preexponential factor remained unchanged. Since all the solid solutions plotted in the figure had a composition of x ¼ 0.5, carrier concentrations will almost be the same, which was also supported by the constant preexponential factor. On the contrary, the activation energy and the ionic conductivity are correlated with the cell volume. The increasing cell volume will open the conduction channel for the Liþ ions and make the Liþ ions more mobile, which lowers the activation energy and increases the conductivity. Many of the oxysalts do not contain any transition metal elements unlike NASICON-type or perovskitetype oxides described below. When the solid electrolyte contains transition metal elements, they are reduced in contact with lithium anodes. In other words, solid electrolytes containing transition metals are not compatible

0 (S cm−1)

Si P

Ge Si P As

Si V

Ge Ge Ti Ti As V As V

106

105

300 K (S cm−1)

104 10−4 10−5 10−6

Ea (kJ mol−1)

10−7 52 48 44 40 330

340

350

3

V (Å )

Figure 1 Specific ionic conductivity at 300 K, s300K; activation energy for conduction, Ea; and preexponential factor, s0, as a function of unit cell volume for x Li4M IVO4  (1 – x)Li3MVO4 system (M IV ¼ Ge, Ti and MV ¼ As, V) systems with x ¼ 0.5.

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with lithium anodes. Therefore, in spite of the low conductivities, g-Li3PO4-type oxysalts have been used in solid-state lithium batteries for a long time. The first lithium battery with the oxysalt electrolyte was already fabricated in thin-film form in 1983. T. Kudo and coworkers assembled the thin-film battery with lithium anode and titanium sulfide (TiS2) cathode. The electrolyte used in the battery was Li3.6Si0.6P0.4O4, which was used in the battery in its amorphous state. Although the specific conductivity at 25 1C was only 5  106 S cm1, thin-film formation shortened the diffusion length in the electrolyte layer, and the battery was operatable at a current density of 16 mA cm2. Though g-Li3PO4-type oxysalts including LISICON are available only in thin-film batteries, they have made a big contribution to recent research on solid-state lithium batteries; it has evolved into two important materials: Lipon and thio-LISICON. J. B. Bates and coworkers found when lithium silicates, lithium phosphates, or lithium phosphosilicates are formed in thin films by radio frequency (r.f.) magnetron sputtering, nitrogen is incorporated into the films and enhances the ionic conductivity. When lithium phosphate is sputtered in pure nitrogen atmosphere, its specific ionic conductivity increases to 2  106 S cm1 (25 1C) in its amorphous phase, which is called Lipon. Although the Lipon itself is amorphous, the crystalline counterpart has g-Li3PO4 structure, and thus it can be categorized into g-Li3PO4-type oxysalts. Many types of thin-film batteries have been reported using Lipon as the electrolyte. Because of the high stability of Lipon, they show excellent performance including a long cycle life, and 5 V cathodes are reported to be available in the battery. Since sulfide ions more weakly attract lithium ions than oxide ions do, sulfides should generally have higher conductivities than oxides. It may suggest that a sulfide iso-structure to an oxide electrolyte will show very high ionic conductivity. However, only glasses have been reported. Of course, their specific ionic conductivities were of the order of 103 S cm1, much higher than oxide glasses; there were no crystalline sulfides with fast Liþion conduction. In 2000, R. Kanno and coworkers reported that Li4GeS4 has g-Li3PO4 structure. Although its conductivity is of the order of 107 S cm1 and not very high, the finding led to many lithium-ionic conductors derived from the materials by aliovalent substitution. The sulfides are categorized into thio-LISICON family. The highest conductivity of 2.2  103 S cm1 (25 1C) with an activation energy for conduction of 20 kJ mol1 was achieved at a composition of Li3.25Ge0.25P0.75S4, which is one of the highest conductivities in solid electrolytes. After 5 years, M. Tatsumisago and coworkers found that a precipitated phase from Li2S–P2S5 glasses shows high ionic conductivities. Recent structure

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Secondary Batteries – Lithium Rechargeable Systems – Lithium-Ion | Electrolytes: Solid Oxide

analysis has revealed that the glass ceramics also have thioLISICON-related structure. When the glass ceramics were precipitated from a sulfide glass with a composition of 70Li2S–30P2S5, they show a very high conductivity of 3.2  103 S cm1 at ambient temperature and a remarkably low activation energy of 12 kJ mol1.

NASICON-Type Phosphates The highest ionic conductivities of g-Li3PO4-type oxysalts are of the order of 105 S cm1, though they are good ionic conductors, which are much lower than those of sulfides and nitrides. One of the reasons is high Li–O bonding energy; oxide ions strongly attract lithium ions, leading to lesser mobility. Oxide ions form a closepacked array in g-Li3PO4 structure, and lithium ions should migrate through the narrow conducting channels among the oxide ions. Relation between the cell volume and the conductivity shown in Figure 1 also clearly demonstrates it; wider channel is necessary for higher ionic conductivity. In 1976, Na1þxZr2P3xSixO12 was found to have fast Naþ-ion conduction, which is called NASICON. In NASICON, ZrO6 octahedra are linked by PO4 tetrahedra to form three-dimensional skeleton structure as shown in Figure 2. Since the open structure was considered to fit fast ionic conduction, many attempts have been made to obtain lithium-ion-conducting counterparts. However, simple substitution of sodium with lithium did not give a

good ionic conductor; ionic conductivity of LiZr2(PO4)3 was lower than 109 S cm1. As revealed by systematic studies of J.-M. Winand and coworkers and H. Aono and coworkers, ionic conduction is strongly correlated with the size of the skeleton network. NaM2(PO4)3 shows the highest conductivity at M ¼ Zr. However, since Liþ ion has smaller ionic radius than Naþ ion, the skeleton framework consisting of ZrO6 octahedra is too large for high ionic conductivity. Figure 3 shows the activation energies for conduction in the bulk obtained for various types of NASICON-type phosphates plotted as a function of the cell volume. Although it includes the data for different types of phosphates, which were derived from LiTi2(PO4)3, LiGe2(PO4)3, and LiHf2(PO4)3, activation energies are on a single curve with a minimum at the cell volume of 1310 A˚3. It is very clear from the figure that the ionic conductivity of LiM2(PO4)3 is the highest at M ¼ Ti, which is a smaller tetravalent cation than Ziroconium. Not only the lattice size but also the carrier concentration is an important factor for high ionic conductivity. Carrier concentration is, in general, controlled by doping aliovalent cations. Indeed, when part of the tetravalent titanium was substituted with trivalent cations including aluminum, scandium, yttrium, and lanthanum resulting in the formation of Li1þxAxTi2x(PO4)3 inducing LiM 2(PO4)3 M = Ge

M = Ti

M = Hf

0.45

0.40 ZrO6

Ea (eV)

PO4 0.35

Na

0.30

0.25 1200

1300

1400

1500

V (Å3)

Figure 2 Crystal structure of sodium super ionic conductor (NASICON).

Figure 3 Activation energy for conduction, Ea, of LiTi2(PO4)3, LiGe2(PO4)3, LiSn2(PO4)3,LiHf2(PO4)3, and their derivatives plotted against the unit cell volume, V. Circles, inverted triangles, squares, and triangles indicate Ea for Ti, Ge, Sn, and Hf systems.

Secondary Batteries – Lithium Rechargeable Systems – Lithium-Ion | Electrolytes: Solid Oxide

interstitial ions, the ionic conduction enhanced. The conductivities showed the maxima at x ¼ 0.3, and the highest conductivity of almost 103 S cm1 was achieved for A ¼ Al. However, further studies revealed that the enhancement of the conductivity does not come from the optimization of the carrier density but from the improved sinterability of the pellets used in the conductivity measurement. The substitution increased the ionic conductivity of LiTi2(PO4)3 from 106 to 104 S cm1 regardless of types of the trivalent cations. That is, the conductivity increased, regardless of whether the trivalent cation was larger or smaller than Ti4þ ion, or whether the lattice was expanded or shrunk. In addition, the conductivities enhanced not only by doping the aliovalent cations but also by adding lithium phosphate (Li3PO4), lithium borate (Li3BO3), and lithium oxide (Li2O). On the contrary, many samples prepared through the studies revealed strong correlation between the porosity and the observed conductivity, which suggests that the enhancement originated from the densification of the sintered pellets. Contribution of bulk and grain boundary can be separately evaluated by complex impedance analysis. Because characteristic frequencies for the ionic conduction in the bulk and grain boundary are different, when the conductivity is measured by complex impedance method,

responses from the bulk and grain boundary conductions appear as two semicircles in the Nyquist plot. Deconvoluted semicircle appearing in the higher-frequency region is always attributable to a response from the bulk conduction, while that in the lower frequency region originates from the conduction at the grain boundary. The contribution of the bulk and the grain boundary evaluated by the above method is shown in Figure 4. Activation energies for bulk conduction were 0.38, 0.30, and 0.42 eV for germanium, titanium, and hafnium systems, respectively. They are not changed by the aliovalent substitutions or by doping lithium oxide. On the contrary, activation energies for conduction at the grain boundary decreased with the substitution or the doping and showed minima at x ¼ 0.1–0.3. This result strongly suggests that neither the substitution nor the doping affects the ionic conduction in the bulk. They only enhance the sinterability to decrease the grain boundary resistance. Nuclear magnetic resonance (NMR) also gave the evidence for the independency of the substitution of the bulk conductivity. Ions in solids usually give broad resonance lines in NMR due to the dipole interaction among the nuclei. However, when the ions are mobile, the interaction is averaged to narrow the resonance lines, which is called motional narrowing. Most of the NMR signal comes from the bulk; the ions do not have to cross the grain boundaries to give the resonance, and thus the

0.55

0.55

0.50

Ea (eV)

Ea (eV)

0.50

Hf grain boundary

0.45 Hf bulk

0.40

Ge grain boundary

Ge bulk

0.35

Hf grain boundary

0.45

Ge bulk

Ti grain boundary

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Ti grain boundary

Ge grain boundary

Hf bulk

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Ti bulk

Ti bulk

0.30

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0.0 (a)

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0.2 x in

0.4

Li1+xM III xM IV 2-x(PO4)3

0.0 (b)

0.1

0.2

0.3

0.4

y in LiM2(PO4)3 - y Li2O

Figure 4 Activation energies for conduction of bulk and grain boundary components for sodium super ionic conductor (NASICON)type phosphates. Circles, triangles, and inverted triangles in (a) are for Li1þx Alx Ge2–x(PO4)3, Li1þx AlxTi2–x (PO4)3, and Li1þx FexHf2–x (PO4)3, respectively, and those in (b) are for LiGe2(PO4)3, LiTi2(PO4)3, and LiHf2(PO4)3 doped with Li2O, respectively. Open and closed symbols correspond to bulk and grain boundary contribution, respectively.

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mobility of the ions can be evaluated without the influence of grain boundaries in NMR. When the ionic motion in two kinds of phosphates, LiTi2(PO4)3 and Li1.3Al0.3Ti1.7(PO4)3, was investigated by NMR, they gave almost the same line width in spite of the large difference between the conductivities measured by electric measurement, which also led to a conclusion that Liþ ionic motion in the two phosphates is of a similar degree.

Perovskite-Type Oxides Perovskite-type alkaline-earth titanates (ATiO3, A ¼ Ca, Sr, Ba) had been attracting much interest for their dielectricity and ferroelectricity. J. Brous and coworkers succeeded in substituting the alkaline-earth ions with a trivalent rare earth (lanthum) and monovalent alkali ions (lithium, sodium, and potassium) and found hysteresis and a dielectric anomaly in Li1/2La1/2TiO3. Y. Inaguma and coworkers observed increasing capacity upon heating, large dielectric loss, and dielectric relaxation that originated from ionic conduction, and reported that the ionic conductivity of the system at room temperature can be as high as 103 S cm1. Perovskite oxides can be represented by a chemical formula of Li3xLa2/3xTiO3. Figure 5 shows the basic crystal structure of the perovskite-type titanates. Titanium atoms octahedrally coordinated with oxygen atoms occupy the corner of the cube (B-site), and the center of the cube (A site) is occupied by La3þ ion, Liþ ion, or vacancy. When the La3þ ions, the Liþ ions, and the vacancies are randomly distributed over the A sites, the lattice belongs to cubic symmetry (space group: Pm-3m). It should be noted that this representation is too simplified. Recent studies based on neutron diffraction showed that Liþ ions do not reside at the center of the cubes but occupy the center of the bottlenecks, and maximum entropy method used in one of the studies also revealed the Liþion conduction paths as density distribution of lithium nuclei.

TiO6

ap

La, Li, vacancy

Bottleneck

Figure 5 Basic crystal structure of perovskite-type solid electrolyte.

Although the basic structure is simple, the system shows many polymorphs originated from cation ordering or tilting of the TiO6 octahedra. When Li0.5La0.5TiO3 is synthesized by heating at 1350 1C and then quenching, the lattice is cubic. However, when it is slowly cooled, broad reflections appear, suggesting a superstructure. In the slowly cooled phase, Liþ ions and La3þ ions are somewhat ordered to form alternate lithium-rich and lanthanum-rich planes, which lowers the symmetry from cubic to tetragonal to form ap  ap  2ap superlattice and gives the superstructure reflections, where ap is a perovskite parameter defined as cube root of the perovskite unit. Tilting of the TiO6 octahedra also distorts the cubic lattice and lowers the symmetry. Hexagonal symmetry was reported for Li0.5La0.5TiO3, and diagonal distortion originated from cation ordering along the c-direction and tilting of TiO6 octahedra results in O2ap  O2ap  2ap superstructure. When the lithium content is very small (xo0.08), the unit cell sometimes becomes orthorhombic. It is believed that the conductivity in perovskite-type oxides is mainly governed by two factors: bottleneck size and site percolation. When a Liþ ion migrates from an A site to the neighboring one, it should go through a bottleneck surrounded by four oxygen atoms. When a Liþ ion approaches the bottleneck (3d site in the space group of Pm-3m), the oxygen atoms attract the Liþ ions, which acts as a potential barrier for the conduction. Therefore, perovskite-type oxides also show correlation between ionic conductivity and lattice size as NASICON-type phosphates. When the conductivities of Li0.34LnTiO3 are compared for Ln ¼ La, Pr, Nd, and Sm, a clear relation between the conductivity and the lattice parameter can be found as indicated in Figure 6. The smaller the lattice parameter, the higher the conductivity and the lower the activation energy for conduction. That is, the larger bottleneck makes Liþ ions more mobile. Although the tendency indicated in Figure 6 suggests that the conductivity is increased, when the A site is occupied by larger trivalent cations, La3þ is the largest trivalent cation. In order to further enlarge the bottleneck size, Y. Inaguma and coworkers introduced divalent Sr2þ ions, which are larger than La3þ, to the A site in place of the lanthanides. The resultant [(Li1/2La1/2)1xSrx]TiO3 showed a conductivity of 1.5  103 S cm1 at 300 K, which was slightly higher than that in the (Li,La)TiO3 system. Liþ ions in the perovskite structure migrate through the A sites, which are partially occupied by Liþ ions and La3þ ions, and the residuals are left vacant. Because the La3þ ions are immobile and block the migration of the Liþ ions, La3þ should distribute to leave a conduction path through the system. In other words, a group of neighboring Liþ ions and vacancies should percolate

Secondary Batteries – Lithium Rechargeable Systems – Lithium-Ion | Electrolytes: Solid Oxide Threshold

10−1 0.7

0.6

Nd (300 K) 0.5

10−4

 (S cm−1)

 (S cm−1)

10

Sm (470 K)

10−3

10−4

−3

Ea (eV)

10−2

333

Pr (300 K)

10−5

10−5 10−6

0.4

10−6 La (300 K) 10−7

10−7 0.380

0.382

0.384

0.386

0.388

ap (nm)

0.30

Figure 6 Specific ionic conductivity at 400 K (open circles) and activation energy for conduction (closed circles) plotted against perovskite parameter for Li0.34La0.51TiO2.94, Li0.34Pr0.56TiO3.01, Li0.34Nd0.55TiO3.00, and Li0.38Sm0.52TiO2.97.

0.35

0.40

0.45

Figure 7 Specific ionic conductivity measured for xLiTaO3  (1 – x) SrTiO3 and that calculated by percolation theory. Open and closed circles indicate the conductivity measured with Au and Li electrodes, respectively.

through the system. According to the percolation theory, conductivity, s, is generally represented as spðx  xc Þm

0.50

x in x LiTaO3 − (1−x )SrTiO3

NbO6 or TaO6

½1

where the percolation threshold for a simple cubic lattice, xc, is 0.311 7, and the exponent, m, is 2.0 for all threedimensional lattices. Figure 7 shows the conductivity data for LiTaO3– SrTiO3 system, where part of A sites accommodate Liþ ions and Sr2þ ions, and B sites are occupied by tantalum and titanium atoms. The best fit for the conductivity data was obtained by m ¼ 2.270.2 on the assumption of xc ¼ 0.311 7. This value agreed with the theoretical value within the statistical error, supporting three-dimensional ionic conduction in perovskite-type oxides.

Garnet-Type Oxides Oxide electrolytes that are the most recently developed have garnet-type structure. The first report was published in 2003, where Li5La3M2O12 (M ¼ Nb, Ta) were disclosed. Garnets are orthosilicates with a general formula of AII3BIII 2 (SiO4)3, where A and B cations are coordinated with eight and six oxygen atoms, respectively. In Li5La3M2O12, La and M occupy eight- and six-coordinated sites in the garnet-like structure as shown in Figure 8.

La Li

Figure 8 Crystal structure of garnet-type oxide.

Although the conductivities of the garnet-type oxides are lower by an order than of NASICON-type phosphates and perovskite-type oxides, they show remarkable advantages: small grain boundary resistance and stability against lithium metal. In NASICON-type phosphates and perovskite-type oxides, contribution of grain boundaries to the resistance is much larger than that of the bulks, even when the powdered samples are sintered into pellets at temperatures higher than 1200 1C. On the contrary, the contribution of the grain boundaries is in the same order of or less than that of the bulk in garnet-type Li6ALa2Ta2O12

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Secondary Batteries – Lithium Rechargeable Systems – Lithium-Ion | Electrolytes: Solid Oxide

(A ¼ Sr, Ba), even when the samples were sintered at 900 1C. Another advantage of garnet-type oxides is compatibility with lithium anode. Although NASICON-type phosphates and perovskite-type oxides have high specific ionic conductivity of the order of 103 S cm1, they are unstable to electrochemical reduction. Lithium ions are intercalated into the NASICON structure with the reduction of Ti4þ to Ti3þ at 2.5 V versus Liþ/Li. Perovskite oxides also contain Ti4þ, which is reduced at 1.5 V versus Liþ/Li accompanied by insertion of Liþ ions. On the contrary, Li6BaLa2Ta2O12 was reported to be stable against lithium metal. No reaction was observed, even when it was immersed in molten lithium.

Application of Lithium-Ion-Conductive Oxide Electrolytes to Electrochemical Power Sources Solid electrolytes with g-Li3PO4-type structure have been applied to thin-film batteries as mentioned above. However, these electrolytes are required to be used in largesized lithium batteries, in which safety issue is much more serious. When such bulky lithium batteries are fabricated with oxide solid electrolytes, the biggest difficulty is a large resistance at grain boundaries. Although ionic conductivities of NASICON-type or perovskite-type solid electrolytes are of the order of 103 S cm1, the large grain boundary resistance will lower the power density. Of course, sintering process at high temperatures will reduce the resistance to some extent; however, such heat treatment will induce interdiffusion between the electrode and electrolyte materials. The interdiffusion results in the formation of impurity phases at the interface, which will degrade the battery performance. Although the grain boundary resistance is very high, oxide solid electrolytes are available in their thin-film form. Of course, what are highly requested may not be thin-film batteries; however, such thin-film oxide electrolytes are effective to improve the performance of bulky batteries as interfacial modification layers. The following are some examples of such interfacial modification. Poly(ethylene)oxide (PEO)-based polymer electrolyte is one of the candidates as electrolytes in lithium polymer batteries because of the low glass transition temperature and flexibility, which are intended to apply to the largesized batteries for electric vehicles. However, as it is not stable to electrochemical oxidation, it cannot be combined with 4 V cathodes, for example, lithium cobalt oxide (LiCoO2) and lithium manganese oxide (LiMn2O4). Therefore, 3 V cathodes, for example, vanadium pentaoxide (V2O5) and lithium iron phosphate (LiFePO4), have been used in the batteries. On the contrary, a research

group of Central Research Institute of Electric Power Industry proposed a new type of battery with a PEObased polymer electrolyte using 4 V cathode, LiCoO2. In order to use LiCoO2, they introduced a thin film of an oxide solid electrolyte, Li3PO4, between the LiCoO2 and PEO electrolyte. The oxide electrolyte layer prevents the direct contact of the polymer electrolyte with the LiCoO2 and oxidative decomposition of the PEO-based polymer electrolyte, which was observed as the reduction of the electrode impedance. Moreover, they demonstrated that the technique enables us to use a 5 V cathode, LiNi0.5Mn1.5O4, in the PEO electrolyte. Another good example of the availability of oxide solid electrolytes is an interfacial modification between oxide cathodes and sulfide solid electrolytes. Sulfide solid electrolytes have higher performance in ionic conduction than oxide ones; the ionic conductivities are of the order of about 103 S cm1, and activation energies for conduction are less than 0.2 eV. In addition, grain boundary resistances are extremely low even without the sintering process; one can make a solid electrolyte layer with low resistance only by cold pressing. In contrast to these advantages, sulfide solid electrolytes have a drawback: when they are connected to high-voltage cathodes, a highly resistive layer is formed at the interface. Weak attraction between sulfide ions and Liþ ions is the reason for high ionic conductivity in sulfides. However, when it is contacted to an oxide, or a high-voltage cathode, the oxide ions in the cathodes attract the Liþ ions more strongly than sulfide ions in the sulfide solid electrolyte to form a lithium-deficient layer on the electrolyte side of the interface. Because carrier concentration is decreased, the conductivity is drastically decreased there, bringing about a high resistance. When an oxide solid electrolyte layer is interposed at the interface, it prevents the development of the lithiumdeficient layer and reduces the interfacial resistance. The interfacial modification improved the high-rate capability of the solid-state lithium battery to be comparable to that of commercialized lithium-ion cells.

Concluding Remarks A variety of oxide electrolytes have been developed as shown here, and high ionic conductivities have been achieved in the electrolytes; however, they are applied only to thin-film batteries. Of course, the ionic conductivities of oxide solid electrolytes are somewhat lower than those of sulfides and liquid electrolytes. The highest specific conductivity of liquid systems used in commercialized lithium-ion cells reaches 102 S cm1; however, this value contains the contribution of the anions. When taking into account that transport number for Liþ ions is unity in the oxide solid electrolytes, it can be

Secondary Batteries – Lithium Rechargeable Systems – Lithium-Ion | Electrolytes: Solid Oxide

concluded that the specific conductivity of 103 S cm1 will be high enough for battery application. Though sulfide solid electrolytes have higher ionic conductivities than oxides and may solve the safety issue of lithium-ion batteries, they are not so chemically stable and difficult to handle. There is no doubt that oxide solid electrolytes are preferable for mass production of solidstate batteries to sulfides. The biggest problem that we have to overcome is the large grain boundary resistance. Recently, a new discipline called ‘nanoionics’ was established to deal with anomalous ionic conduction at hetero-interface. Such studies will give us future prospects of oxide solid electrolytes.

Nomenclature Symbols and Units ap Ea V xc l r r0 r300 K

perovskite parameter activation energy for conduction (kJ mol  1, eV) unit cell volume percolation threshold for a simple cubic lattice exponent specific ionic conductivity (S cm  1) preexponential factor in Arrhenius equation for conduction specific ionic conductivity at 300K

Abbreviations and Acronyms LISICON NASICON NMR PEO

lithium superionic conductor sodium super ionic conductor nuclear magnetic resonance poly(ethylene)oxide

See also: Electrolytes: Solid: Oxygen Ions; Solid: Sodium Ions; Primary Batteries – Nonaqueous Systems: Solid-State: Silver–Iodine; Secondary Batteries – Lithium Rechargeable Systems: All-Solid State Battery; Electrolytes: Glass; Electrolytes: Solid Sulfide; Secondary Batteries – Lithium Rechargeable Systems – Lithium-Ion: Inorganic Electrolyte Batteries.

Further Reading Aatiq A, Me´ne´trier M, Croguennec L, Suard E, and Delmas C (2002) On the structure of Li3Ti2(PO4)3. Journal of Materials Chemistry 12: 2971--2978. Aono H, Imanaka N, and Adachi G (1994) High Liþ conducting ceramics. Accounts of Chemical Research 27: 265--270. Aono H, Sugimoto E, Sadaoka Y, Imanaka N, and Adachi G (1989) Ionic conductivity of the lithium titanium phosphate (Li1þx MxTi2x(PO4)3, M ¼ Al, Sc, Y, and La) systems. Journal of the Electrochemical Society 136: 590--591.

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Aono H, Sugimoto E, Sadaoka Y, Imanaka N, and Adachi G (1990) Ionic conductivity and sinterability of lithium titanium phosphate system. Solid State Ionics 40/41: 38--42. Aono H, Sugimoto E, Sadaoka Y, Imanaka N, and Adachi G (1993) The electrical properties of ceramic electrolytes for LiMxTi2x(PO4)3; yLi2O, M ¼ Ge, Sn, Hf, Zr systems. Journal of the Electrochemical Society 140: 1827--1833. Bates JB, Dudney NJ, Gruzalski GR, et al. (1993) Fabrication and characterization of amorphous lithium electrolyte thin films and rechargeable thin-film batteries. Journal of Power Sources 43: 103--110. Belous AG, Novitskaya GN, Polyanetskaya SV, and Gornikov YI (1987) Study of complex oxides with the composition Li2/3xLi3xTiO3. Inorganic Materials 23: 412--415. Bohnke O, Bohnke C, and Fourquet JL (1996) Mechanism of ionic conduction and electrochemical intercalation of lithium into the perovskite lanthanum lithium titanate. Solid State Ionics 91: 21--31. Bruce PG and West AR (1980) Phase diagram of the LISICON, solid electrolyte system, Li4GeO4–Zn2GeO4. Materials Research Bulletin 15: 379--385. Delmas C, Nadiri A, and Soubeyroux JL (1988) The NASICON-type titanium phosphates ATi2(PO4)3 (A ¼ Li, Na) as electrode materials. Solid State Ionics 28–30: 419--423. Goodenough JB, Hong HY-P, and Kafalas JA (1976) Fast Naþ-ion transport in skeleton structure. Materials Research Bulletin 11: 203--220. Hong HY–P (1978) Crystal structure and ionic conductivity of Li14Zn(GeO4)4 and other new Liþ superionic conductors. Materials Research Bulletin 13: 117--127. Inaguma Y, Chen L, Itoh M, and Nakamura T (1994) Candidate compounds with perovskite structure for high lithium ionic conductivity. Solid State Ionics 70/71: 196--202. Inaguma Y, Liquan C, Itoh M, et al. (1993) High ionic conductivity in lithium lanthanum titanate. Solid State Communications 86: 689--693. Inaguma Y, Matsui Y, Shan Y-J, Itho M, and Nakamura T (1995) Lithium ion conductivity in the perovskite-type LiTaO3–SrTiO3 solid solution. Solid State Ionics 79: 91--97. Itoh M, Inaguma Y, Jung W-H, Chen L, and Nakamura T (1994) High lithium ion conductivity in the perovskite-type compounds Ln1/2Li1/2TiO3 (Ln ¼ La, Pr, Nd, Sm). Solid State Ionics 70/71: 203--207. Kanehori K, Matsumoto K, Miyauchi K, and Kudo T (1983) Thin film solid electrolyte and its application to secondary lithium cell. Solid State Ionics 9&10: 1445--1448. Kanno R and Murayama M (2001) Lithium ionic conductor thio-LISICON, the Li2S–GeS2–P2S5 system. Journal of the Electrochemical Society 148: A742--A746. Kanno R, Hata T, Kawamoto Y, and Irie M (2000) Synthesis of a new lithium ionic conductor, thio-LISICON lithium germanium sulfide system. Solid State Ionics 130: 97--104. Khorassani A, Izquierdo G, and West AR (1981) The solid electrolyte system, Li3PO4–Li4SiO4. Materials Research Bulletin 16: 1561--1567. Kobayashi Y, Miyashiro H, Takei K, et al. (2003) 5 V class all-solid-state composite lithium battery with Li3PO4 coated LiNi0.5Mn1.5O4. Journal of the Electrochemical Society 150: A1577--A1582. Kochergina LL, Khakhin NB, Porotnikov NV, and Petrov KI (1984) A physicochemical study of the series (LiLn)1/2TiO3. Russian Journal of Inorganic Chemistry 29: 879--883. Kunugi S, Kyomen T, Inaguma Y, and Itoh M (2002) Investigation of isotope effect of lithium ion conductivity in (La,Li)TiO3 single crystal. Electrochemical and Solid-State Letters 5: A131--A134. Li S-C and Lin Z-X (1983) Phase relationship and ionic conductivity of Li1þxTi2xInxP3O12. Solid State Ionics 9&10: 835--838. Mizuno F, Hayashi A, Tadanaga K, and Tatsumisago M (2005) New, highly ion-conductive crystals precipitated from Li2S–P2S5 glasses. Advanced Materials 17: 918--921. Mizuno F, Hayashi A, Tadanaga K, and Tatsumisago M (2006) High lithium ion conducting glass ceramics in the system Li2S–P2S5. Solid State Ionics 177: 2721--2725.

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Secondary Batteries – Lithium Rechargeable Systems – Lithium-Ion | Electrolytes: Solid Oxide

Murayama M, Sonoyama N, Yamada A, and Kanno R (2004) Material design of new lithium ionic conductor, thio-LISICON, in the Li2S–P2S5 system. Solid State Ionics 170: 173--180. Ohta N, Takada K, Zhang L, Ma R, Osada M, and Sasaki T (2006) Enhancement of the high-rate capability of solid-state lithium batteries by nanoscale interfacial modification. Advanced Materials 18: 2226--2229. Rodger AR, Kuwano J, and West AR (1985) Liþ ion conducting solid solutions in the systems Li4XO4–Li3YO4: X ¼ Si, Ge, Ti; Y ¼ P, As, V; Li4XO4–LiZO2: Z ¼ Al, Ga, Cr and Li4GeO4–Li2CaGeO4. Solid State Ionics 15: 186--198. Seki S, Kobayashi Y, Miyamoto H, Mita Y, and Iwahori T (2005) Fabrication of high-voltage, high-capacity all-solid-state lithium polymer secondary batteries by application of the polymer electrolyte/inorganic electrolyte composite concept. Chemistry of Materials 17: 2041--2045. Shannon RD, Taylor BE, English AD, and Berzins T (1977) New Li solid electrolytes. Electrochimica Acta 22: 783--796. Stramare S, Thangadurai V, and Weppner W (2003) Lithium lanthanum titanates: A review. Chemistry of Materials 15: 3974--3990.

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